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Full text of "Electrical contacts to P-type zinc telluride and gallium nitride"











First I would like to thank my parents, Henry and 
Andrea Trexler, for all their love and support over the 
years. I am especially grateful that over the five years I 
have been in graduate school, they never once asked the 
question "When are you finally going to get out of school?" 

I would also like to thank my graduate advisor Dr. Paul 
Holloway. I don't think I have ever met a man who knows so 
much about so many different things. His knowledge and 
support have been influential towards my maturation as a 
student and researcher. 

I would also like to acknowledge all the people who 
aided in the work presented in this dissertation. At the 
University of Florida, L. Calhoun for the ZnTe samples, E. 
Lambers for help with the AES analysis, M. Puga-Lambers for 
the SIMS results, and a special thanks to my TEM pit crew. 
I would also like to thank M.A. Khan from APA Optics and B. 
Karlicek from Emcore for the GaN samples. Finally, a big 
thanks goes out to M. Mier at WPAFB for all his help with 
the temperature dependent I-V data. 

I appreciate all the help, conversation, and general 
good times over the past five years from all the members of 
the Holloway group, past and present. It has been enjoyable 


working with each and every one of them. 

Last but not least, I thank all the friends I have made 
during my time in Gainesville, Miller, Rock, Fahey, 
Brunstein, Karen, Wish, JV, Brent, and everyone else I may 
not have mentioned, for making life in Gainesville bearable. 
I have many great memories of my time here that I will carry 
with me always. 








Schottky and Ohmic Contacts 6 

Carrier Transport 11 

Interfacial Reactions 15 

II-VI Materials 17 

Growth and Doping 17 

ZnSe epitaxial layer growth 17 

ZnTe epitaxial layer growth 18 

Schottky and Ohmic Contacts 19 

ZnSe 19 

ZnTe 21 

III-V Materials 23 

GaN Epitaxial Layer Growth and Doping 23 

Substrates 23 

Buffer layers 25 

Autodoping 26 

Contacts 28 

n-GaN 28 

p-GaN 32 


Introduction 39 

Deposition and Processing 39 

p-ZnTe Contacts 39 

p-GaN Contacts 41 

Characterization 43 

Electrical Characterization 43 

Surface Composition Analysis 45 

Surface Morphology 46 

Interfacial Reaction Products 46 

4 p-ZnTe CONTACTS 55 

Introduction 55 


Results 56 

I-V Results 56 

AES 57 

SIMS 58 

Surface Morphology 58 

Discussion 59 

Summary 65 

5 p-GaN CONTACTS 80 

Introduction 80 

Results 81 

Au 81 

I-V results 81 

AES 81 

Pd/Au 82 

I-V results 82 

AES 83 

Ni/Au 84 

I-V results 84 

AES 85 

Ni/C/Au 86 

I-V results 86 

AES and SIMS 87 

Cr/Au 87 

I-V results 87 

AES 88 

XTEM and EDS 89 

Temperature Dependent I-V 90 

Discussion 91 

Metal Contact Schemes 91 

Au 92 

Pd/Au 93 

Ni/Au 94 

Ni/C/Au 97 

Cr/Au 97 

Temperature Dependent I-V 104 

Summary 107 


p-ZnTe Contacts 152 

p-GaN Contacts 153 


p-ZnTe Contacts 155 

p-GaN Contacts 155 



Abstract of Dissertation Presented to the Graduate School 

of the University of Florida in Partial Fulfillment of the 

Requirements for the Degree of Doctor of Philosophy 




December, 1997 

Chairperson: Dr. Paul Holloway 

Major Department: Materials Science and Engineering 

The electrical and structural properties of sputter 

deposited metal contacts to p-ZnTe or p-GaN have been 

studied using current -voltage (I-V), Auger electron 

spectroscopy (AES) and secondary ion mass spectrometry 

(SIMS) . Contacts of Au on p-ZnTe were heat treated over a 

range of 150-350°C for times up to 90 minutes. As-deposited 

the contacts had a breakdown voltage of -0.5 V. Heating at 

T*200°C for 15 minutes resulted in ohmic behavior (linear I- 

V curves) due to formation of near-surface acceptors from 

doping by Au without interfacial phase formation. The 

maximum current density was found to be 2.3 A/cm 2 at 5 volts 

following a 250°C heat treatment for 15 minutes. Diffusion 

of Au, which resulted is a highly resistive contact at 

350°C, was extensive and resulted in compound formation. 


For p-GaN, Au contacts were sputter deposited while 
Ni/Au, Ni/C/Au, Pd/Au and Cr/Au contacts were deposited by 
electron beam evaporation. The contacts were heat treated 
from 200-600°C for times up to 30 minutes. Contacts of 
Pd/Au and Cr/Au were also rapid thermally annealed (RTA) at 
900°C for 15 seconds. Analytical data showed that Au did 
not react with GaN up to 400°C, therefore a rectifying 
contact was observed. A Ni layer between the Au and GaN led 
to dissociation of the GaN at 600°C. Increased electrical 
transport by the Ni/Au contacts was attributed to increased 
doping of the GaN near-surface region by interfacial carbon 
contamination, where it was calculated that 2xl0 19 cm" 3 C 
atoms were incorporated during annealing. Pd/Au contacts 
were rectifying as-deposited. Increased conduction was 
attributed to the formation of a Au:Pd solid solution below 
400°C and dissociation of the GaN following a 900°C, 15 
second RTA. Cr/Au contacts resulted in ohmic behavior with 
P c s4.3xl0 _1 Qcm 2 after a 900°C, 15 second RTA. The RTA caused 
Cr to dissociate GaN, allowing Cr to increase the near- 
surface doping density. It was postulated that Cr diffused 
into the near-surface region using an interstitialcy 
mechanism and formed substitutional acceptors on the Ga sub- 
lattice. Temperature dependent I-V measurements showed that 
charge transport from Ni/Au, Pd/Au and Cr/Au contacts have 
components of both thermionic emission and field emission, 

VI 1 

with Cr/Au having the largest contribution from field 
emission transport. 

VI 11 


Presently there is a strong desire to fabricate 
optoelectronic devices which emit light in the blue region 
of the visible spectrum. Light emitting diodes (LED's) and 
laser diodes (LD's) emitting in this spectral region have 
found applications in optical data storage and full color 
displays. The present optical storage industry is dominated 
by AlGaAs/GaAs based devices which operate in the near 
infrared (IR) region. By replacing the infrared devices 
with shorter wavelength blue emitting devices, the beam can 
be focused to half the minimum spot size of the longer 
wavelength devices, yielding the ability to write four times 
as much data in the same surface area [And93] . The growing 
interest in bright full-color displays has also magnified 
the need for blue light emitting devices. Semiconductor 
optical devices routinely operate from the IR to green 
wavelengths. If this range could be extended into the blue 
wavelengths, semiconductor components could then emit and 
detect the three primary colors of the visible spectrum 
which would have a major impact on imaging and graphics 
applications [Str92] . This need for materials which emit in 
the blue region with high brightness and efficiency has 

stimulated significant research, and while the materials and 
devices have been difficult to develop, great advances have 
been made in recent years. 

There are two groups of blue light emitting materials 
that are viable for optoelectronic devices. These are the 
materials systems based on ZnSe or GaN. ZnSe is a direct 
band gap II-VI compound semiconductor with a zincblende 
structure, a 5.688 A lattice parameter [Wea76] , a band gap 
of 2.67 eV [Shi80] and an electron affinity of 4.09 eV 
[Swa66] at room temperature. ZnSe, doped n-type, has been 
grown using Cl doping during molecular beam epitaxial (MBE) 
growth [Ohk87] , and p-type using a N 2 radio frequency (RF) 
discharge during MBE growth [Par90] . This ability to 
produce both n- and p-type doping has led to the fabrication 
of ZnSe based light emitting diodes (LEDs) [Qui93] and laser 
diodes (LDs) [Haa91] which emit in the blue to green 
wavelengths of 490-516 nm. Presently, the longest lifetime 
devices for laser output consist of a ZnCdSe/ZnSSe/ZnMgSSe 
single quantum well separate-confinement heterostructure 
laser diode with a wavelength of 514.7 nm, which operated 
more than 100 hours at room temperature under continuous 
wave (CW) operation [Tan96] . 

The other material used for blue light emitting 
devices, GaN, is a direct band gap III-V compound 
semiconductor commonly grown in the wurtzite crystal 
structure with lattice parameters of a=3.189 A and c=5.185 

A [Maru69] , a band gap of 3.4 eV [Str92] and an electron 
affinity measured to be between 3.3 eV [Nem96] and 4.1 eV 
[Pan74] at room temperature. GaN has been grown by 
metalorganic chemical vapor deposition (MOCVD) and doped n- 
type with Si donors [Nak92a] , and p-type using Mg acceptors 
[Nak92b] . The ability to grow both n- and p-GaN led to the 
fabrication of p-GaN/n-InGaN/n-GaN double heterostructure 
blue LEDs with a peak wavelength of 440 nm and InGaN multi- 
quantum well (MQW) laser diodes capable of continuous wave 
operation at room temperature with a peak wavelength of 
400.23 nm [Nak97] . 

Currently, GaN based LEDs are the only commercially 
available blue light emitting devices. This fact not 
withstanding, there are still many areas which require 
further advancement before high quality, long life 
commercial devices can be manufactured and marketed. One of 
the critical problems facing these laser diodes is the short 
lifetime under either pulsed or CW operation. Poor contact 
performance, particularly in the p-type contact, is a 
contributing factor to poor device lifetimes. 

An ohmic contact is a non-rectifying 
metal/semiconductor contact which exhibits linear current 
vs. voltage (I-V) characteristics and has a negligible 
resistance in comparison to the bulk resistance of the 
semiconductor. Ohmic contacts allow power to be efficiently 

applied to the active region of the device as opposed to 
being consumed in the contacts. 

This work will focus on improving the current contact 
schemes to p-ZnSe and p-GaN, and to give an understanding of 
the processes responsible for the formation of long-lived 
robust ohmic contacts to these materials. Background 
information on ohmic contacts and a discussion of the 
metallization schemes used and their effects of the 
inter facial reactions between metal and semiconductor will 
be presented. 

A review of the literature is presented in chapter two. 
Metal /semiconductor contact theory and processes utilized in 
the formation of ohmic contacts are discussed. Then the 
literature on growth and doping is presented, focussing 
primarily on the formation of p-type material. This is 
followed by the various contact schemes that have been 
investigated for these materials. This information will 
first be presented for the II-VI materials ZnSe and ZnTe and 
then for the III-V material GaN. 

Chapter three presents the experimental procedure 
followed in this study. The deposition and processing 
parameters are first presented. This is followed by a 
review of the characterization techniques used to study the 
electrical properties, surface composition, surface 
morphology, structural properties and interfacial reaction 

The results and discussion of the contact studies are 
presented in chapters four and five. Chapter four presents 
the data on Au/p-ZnTe contacts. Heat treatments are 
described that lowered the breakdown voltage in these 
contacts and led to the formation of ohmic contacts . The 
composition and structure of the contacts were analyzed to 
identify changes induced by the heat treatments, which led 
to the ohmic behavior and the eventual breakdown of the 
ohmic contacts. In chapter five, the data for Au, Pd/Au, 
Ni/Au, Ni/C/Au and Cr/Au on p-GaN are presented with 
emphasis on the interfacial reactions upon heat treatment 
and conduction mechanisms in these contacts. Conclusions 
from these studies are summarized in chapter six and 
suggested future research is discussed in chapter seven. 


Schottkv and Ohmic Contacts 

When a metal and semiconductor are brought into 
electrical contact a potential barrier arises from the 
separation of charges at the metal-semiconductor interface 
such that a high resistance region devoid of mobile carriers 
is created in the semiconductor [Sha84] . Two types of 
contacts arise from this orientation: ohmic contacts, which 
are defined as a metal -semiconductor contact that has a 
negligible contact resistance relative to the bulk or 
spreading resistance of the semiconductor [Sze81] , and 
Schottky contacts which rectify current flow across the 
metal -semiconductor interface. Most metal /semiconductor 
contacts are Schottky in nature. 

Figure 2 . 1A shows metal and p-type semiconductor energy 
bands before contact. In this figure, <* B and <t>, are the 
metal and semiconductor work functions, defined as the 
amount of energy required to raise an electron from the 
Fermi level to the vacuum level; x. is the semiconductor 
electron affinity, the energy difference between the vacuum 
level and the lower edge of the conduction band; and E g is 


the semiconductor bandgap. In Figure 2. IB the metal and p- 
type semiconductor have been brought into electrical contact 
in thermal equilibrium. When « m «t>,, electrons flow from the 
metal into the semiconductor until the Fermi level on the 
two sides are aligned. After reaching the semiconductor, 
the electrons recombine with the majority carrier holes 
giving rise to a space charge layer of ionized acceptors. 
The concentration of holes in the space charge region is 
negligibly small compared to the acceptor concentration. It 
follows, that on the semiconductor side of the contact the 
space charge region consists of a depletion layer whose 
thickness W depends on the concentration of ionized acceptor 
atoms. In this configuration the obstacle to current 
conduction in contact/p-type semiconductor systems is the 
energy barrier at the valence band (<I> B ) which can be defined 

®b=X.+E b -<V (2.1) 

From this equation and Figure 2. IB, it can be seen that for 
non-degenerately doped semiconductors, where x, + E g *<l> s , a non- 
rectifying contact should be observed for a metal work 
function greater than the semiconductor work function. 
This dependence of $ B on <t>„ is known as classical 
Schottky theory and is only observed in predominantly ionic 
semiconductors. For covalently bonded semiconductors 
Bardeen [Bar47] pointed out the importance of localized 
surface states in determining the barrier height. Dangling 

bonds formed due to the covalent nature at the surface give 
rise to localized energy states at the semiconductor surface 
with energy levels lying in the forbidden gap. These 
surface states may be continuously distributed or localized 
in the band gap and modify the charge in the depletion 
region affecting the barrier height. These states are 
characterized by a neutral level <l> . The effect of these 
surface states on the energy levels in a metal /p- type 
semiconductor can be seen in Figure 2.2. When equilibrium 
is reached, holes from the semiconductor adjacent to the 
surface occupy states below <t> and the Fermi level at the 
surface aligns with that in the bulk. The surface then 
becomes positively charged and a depletion layer consisting 
of ionized acceptors is created in the semiconductor region 
near the surface. Because of this dipole formation, a 
potential barrier looking from the surface towards the 
semiconductor is obtained even in the absence of a metal 
contact. When a metal is now brought into contact with the 
semiconductor and equilibrium is reached, the Fermi level in 
the semiconductor must change by an amount equal to the 
contact potential by exchanging charge with the metal. If 
the density of surface states at the semiconductor surface 
is very large then the charge exchange takes place largely 
between the metal and the surface states, and the space 
charge in the semiconductor remains almost unaffected. As a 
result the barrier height becomes independent of the metal 

work function and is given by: 

* B =E g -* . (2.2) 

This barrier height is said to be "pinned" by the surface 
states. With Fermi level pinning playing a large role in 
the choice of contact materials, the ability to predict and 
determine whether this phenomenon will occur is of utmost 
importance . 

In a model applicable to all semiconductors [Bar47] , 
the Schottky barrier height, S> B , can be expressed as: 

4> B =S*(s)x„+«> (s), (2.3) 

where m and s refer to the metal and semiconductor, 
respectively, and <X> (s) represents the contribution of the 
surface states. The interface index S*(s) expressed as: 

SMs)=d* /dx m , (2.4) 

gives the dependence of barrier height on the metal 
electronegativity, and x™ is the metal electron affinity 
which is analogous to the metal work function. For 
covalently bonded materials, <J> is large and nearly constant 
therefore S* is nearly zero and <D B is nearly independent of 
Xm- Ionic semiconductors, on the other hand, have a large 
S* and <D B is expected to increase linearly with Xm> Kurtin 
et al. [Kurt69] have suggested that S* is a function of the 
electronegativity difference between cation and anion, (AX) , 
of compound semiconductors which is a measure of the 
ionicity of the material, and they have plotted these data 
as shown in Figure 2.3. The regime where S=l is for ionic 


bonding, while the tail near S=0 is for covalent bonding. 
This plot can be used for a rough estimate as to whether 
Fermi level pinning will occur; it is not an exact 
determination of bonding character. One example is the case 
of Si0 2 which is strongly covalently bonded but has a AX 

There are also two other types of interfaces prevalent 
in metal/semiconductor contacts. These are reactive 
interfaces and samples with an insulating, non-reactive 
interfacial layer [Sha84] . The majority of work on reactive 
interfaces has dealt with silicides e.g. [Ott80, Fre80, 
Bri78] , and will not be illustrated here. In most metal- 
semiconductor contacts before metal deposition, the 
semiconductor surface is prepared by chemical cleaning and a 
thin insulating oxide layer is invariably left on the 
surface of the semiconductor. When the interfacial layer is 
thin enough (i.e.<20A), the potential drop across it is 
negligibly small compared to that in the semiconductor 
depletion region. Such a thin layer is transparent to the 
electrons as the electrons can tunnel through in either 
direction. Because of this, the barrier height 0> B and the 
contact potential difference V t may be unaffected by the 
presence of a thin interfacial layer [Sha84] . Thicker 
interfacial layers however, may increase <t> B . 

One final effect on the measured barrier height is a 
result of the electric field in the depletion region which 


leads to image force barrier lowering. This effect occurs 
whether or not an inter facial oxide is present, and the 
magnitude of barrier lowering, A<D B , is given by: 

AG> B = [ [q 3 N d /8n 2 e d 2 e.] * (Vj-V) ] w \ (2.5) 

where N d is the carrier concentration, e d is the dielectric 
constant in the depletion region, e s is the dielectric 
constant in the bulk semiconductor (e d =e„) and Vi is the 
built-in potential of the barrier which is related to the 
position of the Fermi level. Since image force lowering of 
the barrier results from the field produced by an electron, 
the measured barrier height is not lowered by those methods 
which do not require movement of the electron over the 
barrier, (e.g. capacitance-voltage method) . 

Carrier Transport 

Current flow in a metal -semiconductor junction 
occurs from charge transport between the metal and 
semiconductor. There are four mechanisms by which this can 
occur in Schottky contacts: a) thermionic emission over the 
barrier; b) tunneling through the barrier; c) carrier 
recombination (or generation) in the depletion region; and 
d) carrier recombination in the neutral region of the 
semiconductor which is equivalent to minority carrier 
injection (Figure 2.4 [Sha84]). In general semiconductor 
current transport described by classic Schottky theory is 
determined by the carrier concentration in the 


semiconductor. At moderate temperatures (300K) , mechanisms 
c and d dominate for very low carrier concentrations (~10 14 
cm" 3 ) and are rarely used to form low resistance ohmic 
contacts. For low carrier concentrations (N D il0 17 cm" 3 ), 
conduction across the metal -semiconductor interface is 
dominated by thermionic emission (TE) over the potential 
barrier, which is the case for Schottky barrier junctions. 
For higher carrier concentrations (N D *10 19 cm" 3 ) , the barrier 
width becomes narrow and conduction may take place by 
tunneling through the thinned barrier, which is also known 
as field emission (FE) transport [Sze85] . Ohmic conduction 
is related to a large tunneling component. There also 
exists a hybrid case known as thermionic field emission 
(TFE) , which is actually the mechanism (b) shown in Figure 
2.4, in which there is enough kinetic energy for the carrier 
to be excited from the ground state to a level at which the 
potential barrier is thin enough for tunneling to occur. 
The conditions under which each conduction mechanism is 
expected to control current flow are fundamental to 
understanding formation of ohmic contacts. 

Models to determine the dominant conduction mechanism 
have been devised for thermionic emission and field 
emission. These models allow evaluation of conduction 
mechanism from the current-voltage characteristics of the 
metal /semiconductor contact. A model for thermionic 
emission over a barrier was presented by Wagner [Wag31] and 


Schottky and Spenke [Sch39] . In reverse bias the current is 
given by: 

I=I [exp(-qV R /kT)-l], (2.6) 


I =SA*T 2 exp[(A* B -* B )/kT], (2.7) 

and A* known as Richardson's constant is given by: 

A*=4nm*qk 2 /h\ (2.8) 

In these equations I is the saturation current, q is the 
electron charge, V„ is the reverse bias, k is Boltzmann's 
constant, S is the diode area, &<& B -® B is the effective 
barrier height with AO B determined from Equation 2.5 for 
image force reduction, m* is the majority charge carrier 
effective mass, and h is Planck's constant. 

Tunneling can occur in a Schottky barrier junction in 
either forward or reverse bias. The probability of 
tunneling from the semiconductor to the metal increases as 
the doping density of the semiconductor increases or as the 
potential barrier decreases, e.g. through proper choice of 
the contact metal for an unpinned Fermi level . As mentioned 
before, when a metal and semiconductor are brought into 
electrical contact, a bending of the semiconductor bands 
occurs due to formation of a region depleted of free 
carriers whose thickness W (depletion width) is given by 
[Sze81] : 

W=[(2e,/qN D )|V i -V|] M , (2.9) 

where N D is the doping density, V"i is the built-in barrier 


as discussed earlier, and V is the applied bias. This 
depletion width is shown in Equation 2.9 to be inversely 
proportional to the square root of the carrier density in 
the semiconductor. If the tunneling component dominates the 
current flow, the tunneling current can be given by the 

J t ~exp(-q<D B /E 00 ) , (2.10) 

where E 00 is a tunneling parameter proportional to (N D )* 
which will be described in detail later in this section. 
Equation 2.10 thus indicates that the tunneling current will 
increase exponentially with (N D )*, and from Equation 2.9 the 
tunneling probability is proportional to exp(«> B /W) . 

At higher temperature a significant number of carriers 
are able to rise to an energy where there is a thinner, 
lower barrier and tunneling can occur above the ground state 
(i.e. bottom of conduction band for electrons) but below the 
top of the barrier. Since the number of electrons decrease 
rapidly with energy above the Fermi level due to electron 
occupation probability, there exists an energy E„ (measured 
relative to the bottom of the conduction band) at which the 
contribution of TFE becomes maximum. For a semiconductor in 
which the Fermi level is not pinned, the magnitude of this 
potential barrier can be reduced for p-type materials by 
using a metal with a large work function, which will align 
the Fermi level as closely as possible to the semiconductor 
valence level . 


Tunneling through a Schottky barrier has been analyzed 
theoretically by Padovani and Stratton [Pad66] and Crowell 
and Rideout [Cro69] with the following relationships being 

I=I s exp(-qV R /E ) , (2.11) 



E =E 00 coth(E 00 /kT) , (2.12) 

E 00 =(qh/4n) (N d /m*e,)*, (2.13) 

where the symbols have their usual meanings and the pre- 
exponential factor I, is only weakly dependent on voltage 
and is a complicated function of barrier height, properties 
of the semiconductor, and the temperature. The energy E 00 
is an important parameter in tunneling and kT/E 00 is a 
measure of the relative importance of TE versus FE [Pad71] . 
At low temperatures E 00 may become large compared to kT, 
then E =E 00 and field emission dominates. At high 
temperatures where E 00 «kT, then E =kT and thermionic emission 
dominates. For intermediate values of temperature, E 00 =kT 
and thermionic field emission transport dominates current 
transport . 

Interfacial Reactions 

Interfacial reactions may also play a large role in the 
formation of ohmic contacts. A review by Kim et al . 
[KimTJ97] demonstrated the role of interfacial reactions for 


ohmic contacts to wide bandgap semiconductors . Heat 
treatments are often used to form tunneling contacts by 
alloying between the contact metal and the semiconductor 
[Hal50] . The simplistic idea is to choose a metal or metals 
(elements) which create a highly doped semiconductor region 
and leads to ohmic behavior through tunneling transports. 
In other cases, heat treatments result in interfacial 
compounds, such as metal/silicide/silicon contacts. In 
these contacts, the metal/semiconductor interface is 
eliminated and replaced by two new interfaces, a 
metal /compound and a compound/ semiconductor interface. 

The most widespread use of interfacial reactions to 
form ohmic contacts has been in GaAs . The consequences of 
these interfacial reactions vary, but dissociation of the 
compound semiconductor is a necessary, but not sufficient 
condition for formation of good contacts. In those cases 
where decomposition results in an ohmic contact, the surface 
doping concentration in generally increased [Hoi 97] . The 
proposed mechanism for this incorporation of dopant is an 
epitaxial regrowth of the GaAs with the regrown layer having 
an increased dopant concentration. Sands et al . [San88] 
documented epitaxial regrowth of GaAs at the reaction 
interface for the specific case of Si /Ni /GaAs. Through a 
series of heat treatments, (250°C for 30 minutes followed by 
350°C for 30 minutes) they detected a regrown GaAs region 
using XTEM. Holloway et al. [Hol91] showed that interface 


segregation of Si dopants may also occur in the region of 
decomposed GaAs. This was the first direct proof that 
segregation of dopant from the semiconductor (as opposed to 
lattice incorporation from the metallization) could lead to 
ohmic contact formation. 

II-VI Materials 

Growth and Doping 

ZnSe epitaxial laver growth 

II-VI compound semiconductors such as ZnSe (E g =2.67 eV 
at RT) have long been promising materials for the 
fabrication of efficient blue light emitting diodes [Qiu91] . 
One of the objectives of this work is to understand the 
factors in the formation of ohmic contacts to p-doped II-VI 
wide bandgap semiconductors, so growth and doping of these 
p-type materials will be discussed in this section. 

Initial efforts to incorporate a substitutional 
acceptor impurity into ZnSe during crystal growth focused on 
lithium doping during molecular beam epitaxial (MBE) growth 
[Dep89] . Lithium was shown to have a maximum net acceptor 
density of lxlO 17 cm" 3 , above which strong compensation 
occurred resulting in highly resistive ZnSe. In addition, 
Li is mobile in ZnSe and doping was non-reproducible. 
Oxygen has also been reported as a non-reproducible dopant 


in ZnSe layers grown by MBE, [Aki89] with the largest 
reported net acceptor density of 1.2xl0 16 cm" 3 . In 1990, 
Park et al . [Par90] introduced a novel doping technique 
using nitrogen atom beam doping from an RF discharge source 
during MBE growth. Repeatable net acceptor concentrations 
as large as 3xl0 17 cm" 3 were reported using this technique. 
More recently a maximum value of N A -N D =1.8xl0 18 cm" 3 was 
reported by Ohtsuka et al . using an electron cyclotron 
resonance (ECR) plasma source [Oht93] . 
ZnTe epitaxial laver growth 

The wide-bandgap semiconductor ZnTe (E =2.3 eV) is a 
II -VI material which has been used to form ohmic contacts to 
p-ZnSe [Hie93] . Although bulk crystal ZnTe is naturally 
doped p-type by an intrinsic defect, considerable difficulty 
has been experienced in achieving reproducible, high quality 
p-type ZnTe using MBE. Such difficulty has generally been 
associated with the physical incorporation of the dopant 
species due to a low sticking coefficient [Han93] . For 
example, Kitagawa et al . [Kit81] investigated antimony 
doping at flux levels greater than those of the host 
elements, with the F Te =2F Zn . Doping levels of lxlO 18 cm" 3 were 
reported for ZnTe homoepitaxially grown on ZnTe substrates. 
Hishida et al . [His89] used elemental phosphorus to obtain 
p-doping of ZnTe grown by MBE on GaAs substrates. They 
found that the best films were grown with a Te/Zn beam 
pressure ratio (BPR) of 4 . A maximum doping of 4xl0 17 cm" 3 


was obtained with phosphorus beam equivalent pressure 
comparable to that of Te. However, the crystalline quality 
deteriorated with this doping, as shown by a broad x-ray 
rocking curve. Han et al . [Han93] succeeded in obtaining 
doping levels exceeding 10 19 cm" 3 in ZnTe by employing the 
same types of nitrogen sources used to dope ZnSe. The 
crystalline quality of ZnTe was maintained. 

Schottkv and Ohmic Contacts 


Many successful ohmic contact schemes to n-ZnSe have 
been reported and have been reviewed by Fijol and Holloway. 
[Fij96a] An important factor for ohmic contacts is Fermi 
level pinning, and experimental data has been collected 
indicating that Fermi level pinning does not occur with ZnSe 
[Che94, Xu88, Hol97] . Thus ohmic contacts should be 
predicted by the metal work function as long as interfacial 
contamination and oxide layers do not interfere with charge 
transport between the metal and ZnSe. However, with the 
electron affinity of ZnSe equal to 4.09 eV, and an E g =2.67 
eV, (E g +x=6.76 eV) , finding a metal with a large enough work 
function is impossible. Still, many metallization schemes 
have been studied for ohmic contacts to p-ZnSe. The most 
basic schemes consisted of non-graded single metal contacts. 
Fijol et al. [Fij95] investigated the effects of heat 
treatments on Au and Ag contacts to p-ZnSe and concluded 


that both metals formed rectifying contacts with minimum 
breakdown voltages of 3.0 and 2.3 eV, respectively. The 
lower breakdown voltage for Ag was attributed to oxygen 
doping of the near-surface region. Pseudo-ohmic behavior 
for Au contacts was reported by Akimoto et al . [Aki91] by 
using an oxygen plasma to grow an interfacial oxide before 
metal deposition. They proposed that doping occurred 
leading to the pseudo-ohmic behavior. Chen et al . [Che94] 
reported a decrease in the breakdown voltage of about 0.25 
eV upon deposition of 2-3 monolayers of Se at the interface 
between the p-ZnSe and Au. All of these contacts were 
rectifying in nature. 

A second contact scheme involved the semimetal HgSe. 
Lansari et al . [Lan92] used an in-situ MBE epitaxial layer 
of HgSe to reduce the semiconductor interfacial energy 
barrier to about 0.6 eV, and Fijol et al. [Fij96b] used an 
ex-si tu process of reacting Hg with an in-situ Se capping 
layer to form HgSe. The potential barrier of this contact 
to p-ZnSe was measured to be ~0.55 eV, and it remained 

But the most successful efforts at ohmic contacts to p- 
ZnSe have incorporated ZnTe into the contact scheme. Fan et 
al. [Fan93] used a pseudograded Zn(Te,Se) structure 
sandwiched between a p-ZnTe top layer and a p-ZnSe bottom 
layer. The pseudograded contact region consists of 17 
cells, each of 20A thickness. In each cell both the 


thickness of the ZnTe and ZnSe layers were varied. The 
first cell next to the p-ZnSe layer contained 18A p-ZnSe 
and 2A of p-ZnTe, the next cell 17A p-ZnSe and 3A p-ZnTe, 
and so on. It was proposed that this pseudograded band gap 
region allowed injection of holes from the heavily doped 
ZnTe into ZnSe. A specific contact resistance (p c ) of 2x1 0" 3 
Qcm 2 was measured for these contacts with N A =9.5xl0 17 cm" 3 . 
Hiei et al . [Hie93] employed p*-ZnTe/ZnSe quantum wells 
whose sub-bands were aligned in energy to allow resonant 
tunneling of holes through the multi-quantum well region. 
The ZnSe barriers were 20A thick and the ZnTe wells varied 
from 3 to 17A in thickness. This scheme also included a p- 
ZnTe capping layer. They recorded a p c as low as 5xl0" 2 Qcm 2 
for ZnSe with a hole concentration of 7xl0 16 cm" 3 . Fijol and 
Holloway studied the thermal stability of this contact 
scheme and found them to be thermally unstable. They 
deteriorated badly under modest heating, or with high 
current densities (*1000 A/cm 2 ) [Fij96] . 

When graded or MQW contacts of ZnTe /ZnSe are used for 
ohmic contacts to p-ZnSe, an ohmic contact is required to p- 
ZnTe. For this contact both single and multi-layer metal 
contacts have been studied. Ozawa et al . [Oza93] 
investigated Au, Au/Pt, Au/Pd, and Au/Pt/Pd contacts to p- 
ZnTe and reported the lowest specific contact resistance of 
p c =4.8xl0" 6 Qcm 2 for Au/Pd annealed at 200°C. They proposed 


that a reaction took place between Pd and ZnTe and assumed 
that this reaction was more likely to take place at elevated 
temperatures. The exact role of the Pd layer in the 
formation of the ohmic contacts was not determined. The 
contacts were not thermally stable, with an increase in 
specific contact resistance occurring after a 250°C, 3 
minute anneal for all contacts, (e.g. an increase to 
p c =2xl0~ 4 Qcm 2 for Au/Pd, after 200°C, 3 minutes) . Mochizuki 
et al. [Moc94] investigated Au/Pt/Ti/Ni and proposed that 
the reaction between Ni and ZnTe lowered the specific 
contact resistance (p c =7xl0~ 6 Qcm 2 as-deposited) . Insertion 
of the Ti layer was reported to improve the thermal 
stability up to 350°C. In later work, this group detected 
formation of NiTe 2 from the reaction between Ni and ZnTe 
upon annealing at 300 °C for 3 minutes, which apparently 
played an important role in the lowering of the specific 
contact resistance of Au/Pt/Ni/Ti [Moc95] . Finally, Ohtsuka 
[Oht95] investigated Au/p-ZnTe contacts using an oxygen 
plasma cleaning and HCl treatment before Au deposition. A 
Te-rich layer was formed on the surface. Specific contact 
resistances as low as 5.8xl0" 6 Qcm 2 were reported on p-ZnTe 
(N A =2xl0 18 cm" 3 ) . There was no mention of any reaction of Au 
with the underlying ZnTe. 


III-V Materials 

GaN Epitaxial Laver Growt h and Doping 

As previously mentioned, a second objective of this 
research was to understand the formation of ohmic contacts 
to p-GaN. The III-V nitrides, in particular GaN, are also 
promising materials for semiconductor device applications in 
the blue and ultra-violet (UV) wavelength regions [Str93] . 
inN, GaN, and AlN with direct band gaps of 1.9 eV, 3.4 eV, 
and 6.2 eV, respectively [Str93] , can be grown in the 
wurtzite crystal structure to form a continuous alloy system 
with the ability to grade the wavelength of the emitted 
light over the entire visible spectrum. 

However, two intrinsic material problems have inhibited 
the growth of high quality p-GaN. First, no lattice matched 
substrate for GaN has been found so all films have had a 
large defect density (~10 10 " 12 cm" 2 ) . Second, GaN 
intrinsically dopes n-type. The methods and treatments 
discovered to overcome these difficulties and resulting 
production of high quality p-GaN will be discussed below. 

Despite the large lattice mismatch (-14%) and the 
difference in thermal expansion coefficients, the majority 
of GaN growth studies have used the basal plane (0001) of 
sapphire (Al 2 3 ) [Str92] . Other substrates investigated 
with limited success have included SiC [Kim94 and Bot95], 


ZnO [Sit90] , and MgO [Pow90] , all of which are zinc-blende 
crystal structure. More recently Kuramata et al. [Kura95] 
and George et al . [Geo96] have grown GaN on the (111) and 
(100) planes of magnesium aluminate (MgAl 2 OJ which has a 
spinel crystal structure and a lattice mismatch of -10% with 
GaN. One of the advantages of this substrate and Sic is the 
ability to cleave (not possible with sapphire) which 
provides the possibility of facets to be used as a cavity 
mirror in a laser diode. 

Lithium aluminate (LiAl0 2 ) and lithium gallate (LiGa0 2 ) 
single crystals are also attracting attention as substrates 
for epitaxial growth of GaN [Nic96] . The small mismatches 
with GaN (-1.5% for LiAl0 2 and 0.2% for LiGa0 2 ) should lead 
to lower defect concentrations [Wil96] . However, the 
crystal quality of MBE grown GaN films on LiGa0 2 , as 
determined by photoluminescence (PL) measurements, was 
similar to films grown on sapphire. This was attributed to 
the presence of strain due to the difference in thermal 
expansions coefficients between layer and substrate [And97] . 
More recent work has occurred on GaN layers grown 
homoepitaxially on bulk single crystal GaN using MBE [Gas96] 
and metalorganic chemical vapor deposition (MOCVD) [Pon96] . 
While homoepitaxial approaches are promising, the quality of 
epilayer GaN has not improved dramatically (p x ~10 7 -10 8 cm" 2 ) . 


Buffer layers 

Sapphire is still the primary substrate for GaN growth, 
therefore all research groups have used low temperature 
buffer layers between the substrate and GaN film. The 
purpose of these buffer layers is to reduce the detrimental 
effects of the lattice mismatch by trapping defects in the 
layer and not allowing them to propagate to the deposited 
film. Amano et al . [Ama90] used a 500A thick AlN buffer 
layer grown at 600°C and Khan et al . [Kha92] used a 250A 
thick AlN layer grown at 550°C. Nakamura et al. [Nak94] 
have incorporated a 250A GaN buffer layer grown at 500°C. 
In all these cases, the buffer layers were grown by MOCVD. 
More recently, Li et al. [Li95] have reported that a double 
buffer layer with GaN and AlN as the first and second layers 
leads to mirror-like films across the entire substrate. 
Kuznia et al. [Kuz93] investigated the crystalline quality 
of both GaN and AlN buffer layers grown at 500°C. They 
determined, using low-energy electron diffraction (LEED) and 
cross-sectional transmission electron microscopy (XTEM) , 
that the AlN buffer layer was amorphous upon deposition and 
become single crystal following deposition of the GaN 
epitaxial film at 1000°C. Under similar conditions, the GaN 
buffer layer was also amorphous before GaN epi layer growth, 
and then appeared to convert mostly to single crystal form 
with some buried polycrystalline or amorphous regions. 



As pointed out earlier, unintentionally doped GaN has 
in all cases been observed to be n-type [Str92] . There is 
no consensus on the cause of this autodoping, but many 
believe nitrogen vacancies are the intrinsic defect 
responsible for n-type conductivity [Pan73a] . By proper 
growth, the electron concentrations of GaN doped with Si 
were in the range of 10 17 -2xl0 19 cm" 3 [Nak92a] . 

Many potential p-dopants have been investigated, 
including Zn and Mg [Maru69] and Be and Li [Pan73b] , with 
all showing highly compensated material in the absence of 
post-deposition treatments. In addition to compensation of 
donors resulting in semi-insulating layers, low activation 
of the acceptor species was also observed. One of the 
causes of this low activation is the relatively deep level 
of acceptors in GaN. For example, the ionization level of 
Mg in GaN is -150-165 meV [Aka91] . This limits the 
electrically active acceptors at room temperature to less 
than 1% of substitutional Mg. This leads to the discrepancy 
between dopant concentration and free hole concentration. 
Currently using Mg, in order to obtain lxlO 17 cm" 3 free holes, 
there would need to be lxl 19 cm" 3 Mg atoms in the GaN. At 
this time, there is no shallow acceptor known for GaN to 
overcome this problem. 

Along with the relatively deep acceptor levels in p- 
GaN, low activation was also attributed to introduction of 


hydrogen during the growth of GaN, leading to an in-situ 
hydrogenation process which caused the formation of 
acceptor-H neutral complexes and compensation [Nak92b] . 
This conclusion was supported by Brandt et al . [Bran94] who 
investigated intentional hydrogenation of acceptors and 
found a reduction in hole concentration due to the formation 
of H-acceptor complexes. Pearton et al. [Pea96a] studied 
the role of atomic hydrogen introduced during processing in 
passivating the electrical activity of Mg and C acceptors 
using secondary ion mass spectrometry (SIMS) . They 
determined that H had a diffusivity of >10" u cmV 1 at 170°C 
and passivated the Mg and C acceptors. They found that 
annealing at 450-500°C restored the electrical activity, but 
the hydrogen did not physically leave the films until much 
higher (~800°C) temperatures. In another study with H ion 
implanted samples, they determined that the dissociation of 
Mg-H complexes and the loss of hydrogen from GaN were 
sequential processes. Reactivation occurred at s700°C 
annealing under N 2 , while significant concentrations of 
hydrogen remained in the crystal even at 900°C [Pea96b] . In 
another approach to re-activate the dopant, Amano et al. 
used a low-energy electron beam irradiation (LEEBI) process 
with the accelerating voltage far below the threshold energy 
for atom displacement. They achieved a hole concentration 
of 2xl0 16 cm" 3 [Ama89] . The LEEBI treatment was limited to 
the penetration depth of the electron beam, therefore only a 


thin surface layer was converted to p-type material. On the 
other hand, N 2 annealing re-activated the dopant throughout 
the entire epi layer. 

Attempts to limit the presence of H in the growth 
environment have also been reported. Abernathy et al . 
obtained a hole concentration of 3x1 17 cm" 3 , doping with C 
introduced as CC1 4 using a He carrier gas [Abe95] . Zolper 
et al. obtained p-GaN using ion-implantation of Ca acceptors 
requiring an RTA treatment at 1100°C [Zol96a] , while Brandt 
et al. have reported hole concentrations of up to 10 19 cm" 3 , 
with doping efficiencies of up to 10% at room temperature, 
grown by MBE with no post-deposition treatment [Bran94] . 


With n-type material being much easier to grow than p- 
GaN, the large majority of ohmic contact studies are focused 
on n-GaN. In this section, advances made in contacting n- 
GaN are discussed and related to contact schemes for p-GaN. 
The contacts and heat treatments used to obtain low contact 
resistances are described as well as work investigating the 
interfacial reactions which lead to ohmic contacts. Finally, 
previously reported contact schemes for p-GaN will be 

A major topic of discussion for GaN is whether the 
Fermi level is pinned as for GaAs or is unpinned as for 


ZnSe. A pinned E F would lead to barrier heights not 
dependent on the metal work function. If E F is not pinned, 
small work function metals should give lower potential 
barriers and form ohmic contacts to n-type materials, since 
the Schottky contact theory predicts that the barrier height 
is equal to: 

* B =«»„-Xs. (2.13) 

Recent studies have reported that the barrier height 
was not solely dependent on the metal work function. Guo et 
al. used capacitance-voltage (C-V) measurements to 
determine that Schottky barrier heights of Pt and Pd on n- 
GaN were 1.04 and 0.94 eV [Guo95] . They proposed that since 
the difference in work function between Pt and Pd was 
greater than the difference in measured Schottky barrier 
height, that <&„ did not uniquely determine « B . Wang et al. 
measured Pt and Pd Schottky diodes on GaN and reported no 
dependence of the barrier height on the difference between 
the metal work functions [Wan96] . 

Other work has also shown that the Fermi level is not 
pinned. Foresi and Moustakas deposited Al and Au contacts, 
and found Al to be ohmic as deposited while Au required a 
575°C anneal to obtain ohmic characteristics. They 
speculated that Au ohmic contacts were due to Au diffusion 
into the GaN, but showed no evidence to support this 
speculation. This work supported the idea that the Fermi 
level was not pinned since * A i < *xu (4.08 eV to 5.1 eV) and 


^a^Xgsn (4-1 eV) [For93] . Miller and Holloway investigated 
Ag, Au, and Ti contacts and showed that as deposited Ag 
(* Ag =4.3 eV) contacts were weakly rectifying while Au 
provided strongly rectifying contacts. Ti (* Ti =4.4 eV) was 
rectifying as-deposited but became ohmic upon annealing due 
to the formation of a TiN layer [Mil96] . Finally, Fan et 
al. employed a Ti/Al/Ni/Au contact scheme with an in-situ 
reactive ion etching (RIE) before deposition to remove the 
surface oxide. They obtained a low specific contact 
resistance (p c =8.9xl0~ 8 Qcm 2 ) and speculated that the surface 
defect density following the RIE was not large enough to pin 
the Fermi level [Fan96] . 

Thus the data show that either the Fermi level is 
completely or only partially pinned. It is generally 
speculated that partial Fermi level pinning is due to 
interfacial contamination at the GaN surface rather than an 
intrinsic property of the semiconductor itself. 

The majority of these studies concentrated on 
determining the mechanism of ohmic contact formation. Lin 
et al. grew an InN/GaN short period superlattice (SPS) 
sandwiched between a GaN channel and an inN capping layer. 
This SPS allowed greater tunneling to occur through sub-band 
states which led to specific contact resistances as low as 
6xl0" 5 Qcm 2 , with GaN doped ~5xl0 18 cm" 3 without any post- 
annealing [Lin94a] . Another mechanism for ohmic contact 
formation has been to increase the probability of tunneling 


by increasing the doping concentration of the near surface 
region, as discussed above. Since N vacancies act as donors 
in GaN [Jen89] , increasing their concentrations either by- 
high temperature anneals or interfacial reactions have been 
used. Lester et al . deposited non-alloyed Ti/Al contacts on 
Si implanted GaN where the ohmic character (p c =lxl0~ 5 Qcm 2 ) 
was believed to be due to the 1120°C implant activation 
anneal which generated nitrogen vacancies [Les96] , Zolper 
et al. have collected Auger electron spectroscopy (AES) data 
to support the hypothesis that an AlN encapsulant will 
reduce N loss from the GaN substrate during an anneal at 
1100°C. Without a capping layer, N loss may create an n* 
region at the surface [Zol96b] . Cole et al. investigated W 
contacts on n-GaN and concluded that upon annealing at up to 
1000°C, formation of 3-W 2 N and W-N interfacial phases were 
responsible for the ohmic behavior (p c =8xl0~ 5 Qcm 2 for 
N D =1.5xlO" 19 cm" 3 ). They proposed that the N to form these 
phases out-diffused from the GaN without changing the 
original structure, leading to an accumulation of N 
vacancies near the GaN surface [Col96] . A similar result 
was reported for Ti/Al contacts, in which formation of a 
TiN x interfacial phase was postulated to create N vacancies 
[Bin94] . Lin et al . have speculated that only 1-2 
monolayers of TiN are needed to form N vacancies which would 
generate a 100A layer of GaN with an electron density of 
10 20 cm" 3 [Lin94b] . Ruvimov et al . detected a TiN layer by 


XTEM at the Ti/GaN interface after annealing at 900 °C 
[Ruv96] . Finally, Smith et al. investigated Al contacts and 
found the formation of a thin AlN layer on the surface led 
to increased carrier conduction, presumably due to increased 
N vacancies [Smi96] . 

Various other studies of multilayer contact schemes 
leading to either extrinsic doping of the near surface 
region or phase formation led to decreased contact 
resistances. Miller and Holloway postulated that in 
Au/Si/Ti contacts, Si diffused to the GaN surface and served 
as an n-type dopant [Mil96] . For Ptln 2 annealed at 800°C 
for 1 minute in a high purity Ar atmosphere. In was 
incorporated into the GaN lattice producing (IrixGa^jN and 
Pt(In,Ga) 2 . Formation of (In^a^jN resulted in ohmic 
behavior with p c slxl0~ 3 Qcm 2 [Ing97] . Ping et al . showed 
that annealed Pd/Al/n-GaN contacts formed a Pd 2 Al 3 phase 
with p c =1.2xl0" 5 Qcm 2 upon annealing at 650°C for 30 seconds 
[Pin96] . Luther et al . reported that Ti in Ti/Al contacts 
reduced the native inter facial oxides, Al diffused through 
the Ti to the interface, and an Al-Ti intermetallic formed 
which contacted the GaN. These contacts became ohmic upon 
annealing at 400°C for 3 minutes in an Ar/4%H 2 atmosphere 
(p c =7xl0~ 6 Qcm 2 for N D =5xl0 17 cm" 3 ) [Lut97] . 

In contrast to n-GaN contacts, which have been widely 
studied, there is very little information on p-GaN contacts. 


A limited number of studies have reported Schottky barrier 
heights and that inter facial contamination has limited 
contact formation. However, most reports simply mention the 
p-contact used in a particular device. There has been very 
limited information on the formation of interfacial phases 
or conduction mechanisms in p-GaN. Due to the large E g (3.4 
eV) and x. (4.1 eV) of GaN, a metal would need O m *7.5 eV to 
achieve an ohmic contact with the Fermi level unpinned. 
Unfortunately there are no metals with work functions larger 
than 5.6 eV, [Sze81c] thus interfacial reactions, doping of 
the near surface region, or graded band offsets must be used 
for ohmic contact formation. 

Two groups have reported the effect of metal work 
function on contact properties to p-GaN. Mori et al. 
measured the Schottky barrier heights and a contact 
resistances of Pt, Ni, Au and Ti to p-GaN. Their data 
showed both <X> B and p c decreased with increased metal work 
function, as expected for an unpinned Fermi level [Mor96] . 
Ishikawa et al . investigated the effects of surface cleaning 
on the electrical properties of metal /GaN interfaces, and 
measured p c for Pt, Ni, Pd, Au, Cr, Ti, Al and Ta on p-GaN. 
They showed that removal of the interfacial contamination 
layer between the metal and the GaN did not significantly 
reduce the contact resistance. However, the resistance 
decreased exponentially with increased metal work function. 
Also no reaction between any of these metals and GaN could 


be detected after annealing at temperatures below 500°C. It 
was proposed that for ohmic contacts, a metal is needed 
which would react with GaN during annealing at high 
temperatures [Ish97] . The remainder of information on p-GaN 
contacts consists of the metals used for devices. Early- 
work on p-n junctions and p-GaN/n-InGaN/n-GaN double 
heterostructure blue Light emitting diodes (LEDs) used Au as 
the p-contact [Nak91 and Nak93] . Subsequently, a Ni 
interfacial layer was added between the Au and p-GaN 
[Nak95] . Other groups also used this Ni/Au scheme [Kha95] 
with Molnar et al . reporting thicknesses of 200A Ni/2000A 
Au [Mol95] . All subsequent reports by Nakamura et al . 
regarding their LEDs and laser diodes (LDs) used this Ni/Au 
scheme (e.g. [Nak96]). other groups have used Ni [Sak95] , 
Au + Zn [Kug95] and Ti/Mo/Au (non-ohmic) [Gol93] . None of 
these studies reported the electrical properties of the 
contacts or metallurgical reactions occurring between the 
metals and semiconductor. In summary, while there has been 
a great deal of work on n-GaN contacts, the understanding of 
contacts to p-GaN is very limited and needs to be elevated 
for GaN based blue light emitting devices to reach their 





♦ . 


Figure 2.1A. Metal and p-type semiconductor before contact 













E , 


\ ■* 



^1 1 1 


/ / / / / 

w / 

E v " 

Figure 2. IB. Metal and p-type semiconductor in thermal 
equilibrium after the contact has been made. 






Figure 2.2. Metal/p-type semiconductor energy bands, showing 
the effects of surface states, after contact has been made. 





• SnOj ZnO 


• • 


SiO ? ' 


Go ? Oj 









ZnTa* / 



InSO J^ 

' ' 

1 1 l 














1.6 2.0 2.4 

Figure 2.3. Index of interface behavior, S as a function of 
lattice electronegativity difference AX between the cation and 
anion for various semiconductors [Kur69] . 



E r 


^ fc v 

• Electron 
O Hole 

Figure 2.4. Energy band diagram of a forward biased Schottky 
barrier junction on an n-type semiconductor showing different 
transport processes; (a) thermionic emission, (b) thermionic 
tunneling through the barrier, (c) carrier recombination in 
the depletion region, and (d) hole injection from the metal 
into the semiconductor [Sha84d] . 



The following chapter describes the experimental 
process that was followed to prepare electrical contact 
samples to p-ZnTe and p-GaN. The procedure consisted of 
initial cleaning of the samples followed by metal contact 
deposition. These contacts were then heat treated and 
characterized in terms of their electrical properties, 
surface composition, surface morphology and inter facial 
reaction products . 

Deposition and Processing 

p-ZnTe Contacts 

Investigation of Au contacts was performed on ZnTe/ZnSe 
samples grown on semi-insulating GaAs substrates using a 
custom made Molecular Beam Epitaxy (MBE) system and the 
sample configuration shown in Figure 3.1 [Jeo95] . An 
undoped 0.5 urn ZnSe buffer layer was epitaxial ly deposited 
on the GaAs, followed by a 2 . 3 urn p-ZnSe epilayer. The 
ZnSe/ZnTe multi-quantum well (MQW) structure described by 



Hiei et al. [Hie93] was epitaxially deposited onto the p- 
ZnSe layer. The MQW structure was capped by a HOOA 
heteroepitaxial p-ZnTe layer. An Oxford Applied Research 
Systems nitrogen free radical source was used for the p-type 
dopant layers. The free hole density was measured by a Hall 
experiment to be -3xl0 18 cm" 3 . 

No further surface preparation was done on the as-grown 
ZnTe samples before metal deposition. The Au contacts were 
deposited by sputtering [Ohr92] in a custom built sputter 
deposition system [Tru92] . The system contained two radio 
frequency (RF) and two direct current (DC) 2" diameter 
planar magnetron sputter guns manufactured by US Inc. The 
vacuum chamber was a quartz bell jar with a Viton seal, 
seated on a stainless steel base plate. The system 
consisted of a liquid nitrogen cooled (trapped) diffusion 
pump (Varian VH5-6) which was backed by a Sargent-Welch 
mechanical pump, model #1397. The ultimate pressure reached 
by the system was ~2xl 0" 7 Torr. 

For deposition, the samples were loaded in the sputter 
chamber (Figure 3.2) which was pumped to a base pressure of 
at least 2xl0" 6 Torr. The 1500A thick Au contacts were 
deposited from a 2 inch Pure Tech Inc. sputter target, 
(99.99% purity) using a DC sputter gun with a potential of 
375 V, and a power of -100 W. The gas pressure was 18 mTorr 
argon and a deposition rate of -188 A/minute was achieved. 
To improve the uniformity of the film thickness, the sample 


was rotated about the system axis at 25 rpm. The Au 
contacts were patterned as dot contacts (0.8 mm diameter) 
during deposition using a stainless steel shadow mask. This 
method resulted in electrically isolated discrete Au dot 
contacts separated by -0.4 mm (Figure 3.3). 

Following deposition, the samples were heat treated in 
a quartz tube furnace in flowing forming gas (10% H 2 , 90% 
N 2 ) . Separate Au/ZnTe samples were individually heat 
treated at 150, 200, 250, or 350°C using 15 minute 
increments for total times up to 90 minutes. 

p-GaN Contacts 

Metal contacts were deposited on a variety of GaN 
epitaxial films grown by metalorganic chemical vapor 
deposition (MOCVD) on a polycrystalline GaN buffer layer on 
a (0001) sapphire substrate [Kha91] . Mg acceptors were used 
to obtain p-type conduction and carrier concentrations at 
296K were measured by a Hall technique to be between 5x1 16 - 
4.5xl0 18 cm" 3 . The contact schemes investigated consisted 
of Au, Ni/Au, Ni/C/Au, Pd/Au, and Cr/Au with the first metal 
listed being deposited first and adjacent to the GaN. All 
samples were degreased prior to deposition using acetone, 
methanol, and de-ionized water (DI H 2 0) for 5 minutes 
followed by a N 2 blow dry. Any native oxide was removed 
using a 10:1 DI:HCl etch for 5 minutes followed by a 5 
minute DI rinse and N 2 blow dry. The samples were then 


immediately introduced to the deposition chamber. 

The 2000A thick Au contacts were DC magnetron sputter 
deposited using the previously described sputter system and 
the same target as was used for the Au/p-ZnTe contacts. For 
GaN contacts, the sputter atmosphere was 25 mTorr Ar with a 
power of 40 W. For all GaN contacts, a smaller stainless 
steel shadow mask was used which defined contacts with a 
diameter of -0.5 mm and a spacing of -0.2 mm. 

All contacts other than Au were deposited in a Davis 
Wilder vacuum system with an Airco Temescal four pocket 
electron beam evaporator [Ohr92], powered by a model ES-6 
power supply (Figure 3.4). The system was pumped with a 
liquid nitrogen trapped Varian VH5-6 oil diffusion pump 
backed by an Edwards 40 two stage mechanical pump, providing 
a base pressure of -6xl0" 7 Torr. For all electron beam 
evaporations, the deposition pressure was sl-3xl0~ 6 Torr and 
the charge consisted of metal pellets with the following 
purity's: Ni, (Cerac Inc., 99.95%); Au, (Materials Research 
Corporation, 99.95%); Pd, (AESAR Johnson Matthey Inc., 
99.8%); Cr, (Union Carbide, purity unknown), and C, 
(graphite slug, purity unknown) . The metal layer thickness 
was monitored using a quartz crystal oscillator controlled 
by a Kurt J. Lesker Company QXM-500 controller. The GaN 
contact thickness varied with different experiments and will 
be noted in the results section of Chapter 5. 


Following deposition, all metallization schemes were 
heat treated in the quartz tube furnace previously mentioned 
using either flowing forming gas or N 2 as the ambient. 
Separate samples were heated at 200, 400 or 600°C for 5, 15, 
and 30 minutes. The Ni/Au, Pd/Au and Cr/Au samples were 
also heated to 900°C for 15 seconds in a 50 cm quartz tube 
custom rapid thermal annealing (RTA) furnace with a 25 cm 
hot zone and flowing N 2 as the ambient. 


Both the p-ZnTe and p-GaN samples were characterized 
as-deposited and following each of the previously described 
heat treatments in terms of their electrical properties, and 
also for their surface composition, surface morphology, and 
inter facial reaction products as required. 

Electrical Characterization 

The electrical properties of all contacts were 
investigated using room temperature current-voltage (I-V) 
measurements. The I-V data were obtained by measuring the 
current flow between two adjacent dot contacts under an 
applied bias. An automated system consisting of an IBM PC 
with IEEE-488 communications, a Hewlett-Packard 6112 A DC 
power supply and either a Keithely 488 picoammeter or a 
Hewlett-Packard 3478A multimeter, depending on the current 
range measured, was used to obtain the I-V data. A 


schematic of this setup is shown in Figure 3.5. The 
resulting data, collected over a range of -5 to +5 V, was 
modeled as two back- to-back diode barriers, one forward 
biased and the other reverse biased, and the ohmic or 
rectifying nature of the contacts could be determined by the 
linearity of the I-V curve. The reverse-bias breakdown 
voltage could also be determined from the I-V 
characteristics. Barrier heights were also determined from 
the I-V data, but due to problems involving large ideality 
factors for the contacts, the calculated barrier heights 
were concluded to be inaccurate. 

The charge transport mechanisms for Ni/Au, Pd/Au, and 
Cr/Au contacts were determined using temperature dependent 
I-V measurements. For these measurements, the samples were 
mounted and Au wire bonded onto T05 headers which were 
placed in a liquid nitrogen cooled sample holder. The I-V 
characteristics were determined from -5 to +5 V using a 
Hewlett-Packard 6111A DC power supply with a Hewlett-Packard 
3455A digital voltmeter and the current was measured across 
a lkQ resistor. The temperature ranged from 80-400 K and 
was measured using a Pt resistance thermometer calibrated to 
0.01° controlled by a Lake Shore Cryogenics DRC 82C 
temperature controller. All measurements were taken at a 
pressure of 10 mTorr. 

The p-GaN samples were also characterized in terms of 
their doping character using photoluminescence (PL) [Bru92] 


and Hall effect measurements. The PL data was obtained 
using excitation from a HeCd laser (X=325 nm) at a 
temperature of 22 K for an as-grown sample and also 
following a 900°C, 15 second RTA in N 2 . For the Hall effect 
data, In pads were soldered onto the p-GaN and heat treated 
at 400°C for 3 minutes in N 2 for better adhesion. The data 
was taken with a current of ±0.0001 A and a magnetic field 
of 8 KGs. Following measurement, the In pads were etched 
using a 10:1 DI:HC1 solution and the same sample and 
contacting process was used following the 900°C heat 
treatment . 

Surface Composition Analysis 

Following heat treatment and electrical 
characterization, the elemental surface composition was 
determined by Auger electron spectroscopy (AES) [Hol80] 
using a Perkin-Elmer PHI 660 scanning Auger microprobe 
(SAM), or by secondary ion mass spectrometry (SIMS) [Ben87] 
using a Perkin-Elmer PHI 6600. For contacts to p-ZnTe and 
p-GaN, AES surface survey spectra were recorded over the 
energy range of 50-2050 eV using a 5 keV, 30 nA electron 
beam with a diameter of ~1 um. For depth profiles, a 5 keV 
argon ion beam was used to sputter the sample at a nominal 
rate of -lOOA/min for 6 second intervals after which the 
surface was scanned with the electron beam. SIMS depth 
profiles of Au/p-ZnTe contacts were performed using a 5.5 


keV, 53 nA Cs* ion beam with a 200 um x 200 um raster size 
using a quadruple mass spectrometer to analyze the sputtered 
ions. For p-GaN, SIMS was performed only on the Ni/C/Au 
contacts with a 3 or 5 keV, 30 nA Cs* ion beam. The beam 
had a linear gating of 70% and a 300 um x 300um raster size. 

Surface Morphology 

Changes in surface morphology were investigated using 
optical and scanning electron microscopy (SEM) [Gol92] . All 
samples were viewed following heat treatments with an 
Olympus BH-2 optical microscope with photographic 
capabilities for any large changes in surface morphology. 
SEM was performed on the Au/p-ZnTe samples using a JEOL 6400 
scanning microscope with an accelerating voltage of 15 KV. 

Interfacial Reac tion Products 

For Cr/Au contacts on p-GaN, cross-sectional 
transmission electron microscopy (XTEM) [Edd74] and energy- 
dispersive x-ray spectroscopy (EDS) [Bru92] were used to 
determine interfacial reaction products formed in samples 
subjected to a 900°C, 15 second RTA. These data were 
compared to as-deposited samples. 

For XTEM, the sample preparation technique is extremely 
important in obtaining high quality samples. For these 
cross section experiments, two samples were first glued face 
to face using Gl epoxy (Gatan) . This structure was then 

placed inside a stainless steel notched rod which was in 
turn slid into a stainless steel outer tube with an inner 
diameter just larger than the outer diameter of the rod. Gl 
epoxy was then applied to fix the samples in the notch and 
to fill the area between the rod and the tube. The epoxy 
was cured on a hot plate at 100°C for 30 minutes. After 
curing, samples were cut on a Accutom low speed saw using a 
diamond blade to a thickness of =300 um. Figure 3.6 shows a 
planar schematic view of the samples following the cut. The 
samples were manually polished using SiC polishing paper 
until one side was smooth, flipped over, and polished down 
to a thickness of -100 um. A slotted Cu ring was then glued 
to the sample using Ml epoxy (Micro Measurements Inc.) for 
added stability. Next the samples were flattened to a 
thickness of -30 um on a VCR Group Inc. Dimpler* using a 
tool steel wheel and a 1 mm diameter cylinder. These 
samples were ion milled in a Gatan Dual Ion Mill (Model 600) 
with a voltage of 4-6 kv at 1mA with a final milling angle 
of 12°. After ion milling to produce a hole in the center 
of the structure, the edges of the hole were electron 
transparent. The samples were analyzed by amplitude 
contrast imaging in a JEOL 200CX TEM using a two beam bright 
field condition. Phase contrast imaging was used to obtain 
high resolution images in a JEOL 4000FX high resolution TEM. 
The EDS data was taken with a 400 KeV incident electron beam 
with a spot size of =500A. Characteristic x-rays were 


detected from 1-20 KeV using a Princeton Gamma Tech model 
OS26-J029 detector. 


ZnSe/ZnTe MQW 

Figure 3.1. Schematic diagram of Au/p-ZnTe sample showing 
ZnSe/ZnTe MQW. 


Al Sample Holder 


1 cm 


3 cm 

Ta Wire 


Steel Shield 

2" Au 


Figure 3.2. Schematic of the sample holder and magnetron 
source for the sputter system used to deposit Au conacts. 


Circular Dot Contacts 

d s 

ZnTe 0.8 nun 0.4 mm 
GaN 0.5 mm 0.2 mm 

(A) Top View 

Metal Contacts 

(B) Side View 

Figure 3.3. Top (A) and side (B) views of dot contacts on 
the semiconductor films showing electrical isolation of the 
metal contacts. 








Evaporated Metal 



Figure 3.4. Schematic of electron beam evaporation system. 


Electrical Connections 
IEEE Connections 


Dot Contact 

Figure 3.5. Hardware! setup for current-voltage measurement 



Stainless Steel 

Stainless Steel 

Metal Contact 

GaN Samples 

Figure 3.6. Planar view of Cr/Au on p-GaN sample used for 



Formation and characterization of Au/p-ZnTe ohmic 
contacts are discussed in this chapter. The sputter 
deposited contacts were heat treated in forming gas and 
current- voltage (I-V) measurements were used to determine 
the effect of elevated temperatures on the electrical 
properties. Heat treatments were identified which led to 
the decrease in contact resistance and formation of the 
ohmic contacts. The temperature stability of this contact 
scheme was also studied. Analytical characterization 
techniques, including Auger electron spectroscopy (AES) and 
secondary ion mass spectrometry (SIMS) depth profiling, were 
used to identify elemental surface composition, interfacial 
reactions and compound formation induced by heat treatment. 
These analytical results, along with surface morphological 
changes determined from scanning electron microscopy (SEM) 
were correlated with the electrical results to determine the 
origin of the ohmic behavior and the eventual degradation of 
the contact properties. 




I-V Results 

As-deposited I-V curves were non-linear and 
characteristic of poor back-to-back Schottky contacts 
(Figure 4.1). Upon heating to 150°C for 15 minutes, there 
was a slight reduction in breakdown voltage (V B ) from -0.5 
eV to -0.2 V, but the curves remained slightly non-linear. 
The breakdown voltage was determined by extrapolating the 
linear region of the I-V curve to its x-axis intercept. For 
samples heat treated at 200°C and 250°C for 15 minutes, 
linear I-V curves were measured with decreased contact 
resistance. A maximum current density of 2.3 A/cm 2 at 5 V 
was observed following a 250°C, 15 minute heat treatment. 
Upon heating at 350°C, the I-V curves remained linear but 
the resistance was a factor of 25 higher than for samples 
heated at lower temperatures for equivalent times, 
indicating degradation of the contacts or increased bulk 
resistance of the p-ZnTe (Figure 4.2). Increased time at 
this temperature led to an increased contact resistance. 
The changes that occurred upon heat treatment of the Au/p- 
ZnTe contacts are summarized in Figure 4.3 showing that the 
contacts were thermally stable only for T*250°C. 


AES depth profiles were collected for as-deposited 
samples and after heat treating at 200°, 250°, or 350°C for 
90 minutes. The as-deposited profile (Figure 4.4) shows an 
abrupt Au/p-ZnTe interface indicated by the rapid decrease 
of the Au signal at -22 minutes. For samples heat treated 
at 200° and 250°C (Figures 4.5 and 4.6) the slope of the Au 
signal versus sputter time was not as large as that of the 
as -deposited case. The smaller slopes of the Au and Te 
signals versus depth (at -11 minutes in Fig. 4.5 and -9 
minutes in Fig. 4.6) indicated that Au had begun to diffuse 
into the near surface region of the p-ZnTe leading to a 
slight broadening of the metal/semiconductor interface. 
Even with this limited diffusion, the interface remained 
reasonably planar with no evidence of compound formation. 
In Figure 4.7, the AES depth profile for the 350°C heat 
treatment would be consistent with an interface between Au 
and ZnTe which was no longer planar. Au had diffused 
throughout the ZnTe capping layer and into the underlying 
ZnSe. The change of Au slope leading to a relatively flat 
Au signal from -25-35 minutes of sputtering time suggests 
formation of an inter facial phase or compound. 



SIMS depth profiles were measured on samples as- 
deposited and heat treated at 250° or 350 C C for 90 minutes. 
The as-deposited profile (Figure 4.8) is consistent with a 
very planar Au/ZnTe interface, similar to the AES depth 
profile, with no indication of any Au diffusion into the 
semiconductor. Upon heating at 250°C for 90 minutes (Figure 
4.9) Au had begun to diffuse into the ZnTe. At 5 minutes of 
sputter time the Au signal was larger than the Te signal as 
opposed to data in Figure 4.8 at -4 minutes of sputter time 
where the Au signal was less than the Te signal strength and 
remained lower throughout the profile. At 350°C for 90 
minutes (Figure 4.10), it is readily apparent that Au has 
diffused through the entire ZnTe capping layer and into the 
underlying ZnSe film. 

Surface Morphology 

Changes in the surface morphology, which were 
significant to thermal stability of the Au contacts, were 
observed. The as-deposited Au film was specular reflective 
to the eye and very smooth when examined by scanning 
electron microscopy. After heat treating for 90 minutes at 
T<250°C the Au remained specular to the eye. After 250°C 
for 90 minutes, optical microscopy and SEM showed that the 
ZnTe between the Au dot contacts became hazy (Figure 4.11). 
An AES surface survey showed the presence of Au between the 



dot contacts, creating an extended diffusion zone =100um 
across the ZnTe surface. For the same sample at ti75 
minutes, T=250°C, there was no evidence of this diffusion. 
At 350°C for 90 minutes, a rough surface indicative of 
further degradation was seen (Figure 4.12). 


Ohmic contacts are characterized by their electrical 
properties. This section will discuss three changes in the 
I-V curves upon heat treatment and provide metallurgical 
explanations for the changes. These changes were: 1) 
decreased breakdown voltage at 150°C, 2) increased current 
levels and linear I-V with increased temperature above 
150°C, and 3) dramatic resistance increase and contact 
degradation at 350°C. 

To understand the properties expected from Au/p-ZnTe 
contacts, it must first be determined if the Fermi level is 
pinned. From the Schottky model for a semiconductor with no 
Fermi level pinning, Eq. 2.1 predicts the potential offset 
in the valence band. For Au/p-ZnTe contacts with E g =2.26 eV 
[Mil72], Xzn Te =3.5 eV [Mi 172], and <K Au =5.1 eV, Eq. 2.1 yields 
a value of 0.66 eV. This is in good agreement with the 
experimentally observed as-deposited value of breakdown 
voltage =0.5 eV. This would only be expected if the Fermi 
level was not pinned in p-ZnTe. 


The decreased breakdown voltage upon heat treating to 
150 C C for 15 minutes is attributed to decomposition of the 
thin inter facial contamination layer, evident from the weak 
oxygen signal in Figures 4.4-4.7. The increased current and 
linear I-V curves after heating to 200°C are attributed to 
diffusion of Au into the underlying ZnTe. The analytical 
results from the AES depth profiles (Fig. 4.5 and 4.6) show 
a broadening of the Au/ZnTe interface consistent with Au 
diffusion into the near-surface region of the p-ZnTe. The 
SIMS depth profile for 250°C (Fig. 4.9) also shows direct 
evidence of Au diffusion into the ZnTe capping layer. Au 
has been shown to be principally accommodated 
substitutionally on Zn sites (Au Zn ) and to form a simple 
acceptor level at Ey+272 meV. [Mag80] Thus it is proposed 
that the introduction of Au as an acceptor species in ZnTe 
was responsible for the increased current after elevated 
temperature treatments and therefore the ohmic contact. 

If Au was acting as an acceptor in p-ZnTe and increased 
the doping concentration of the near-surface region, the 
main conduction mechanism in these Au/p-ZnTe contacts should 
have been field emission. To evaluate whether field 
emission or thermionic emission was dominant, thermionic 
emission currents were modeled using Eq. 2.5-2.8. For these 
calculations, the following values were used; <D B for Au on 
p-ZnTe =0.73 eV, T= 300 K, area=0.005 cm 2 (0.8 mm diameter), 
N A =3xl0 18 cm" 3 , hole effective mass=0.60, and low frequency 

dielectric constant e=9.4. [Wag92] The built in potential 
of the junction (VJ was calculated by determining the 
position of the Fermi level and subtracting this value from 
<D B . Based on these calculations, the reverse bias breakdown 
voltage for the Au/p-ZnTe contacts would be 18 V if 
thermionic emission was dominant, rather than the observed 
value of 0.5 eV. Thus thermionic emission was not 
dominating current transport. 

Therefore the properties of the Au/p-ZnTe contacts were 
modeled based on field emission. From the equations for 
field emission, taking the natural logarithm of Eg. 2.11 and 
with some simple rearrangements: 

ln(I)=ln(I s ) + (-qV R /E ) . (4.1) 

Based on this equation, a plot of the In (I) vs. V should 
produce a straight line if tunneling is dominant. Figure 
4.13 shows a plot of In (I) vs. V from 0-5 volts for an as- 
deposited sample as well as samples heat treated at 150, 
200, 250, or 350°C for 15 minutes. As can be seen in all 
cases, the plot of ln(I) vs. V was only linear for V*2V, 
indicating that tunneling was not dominant over the entire 
range of reverse biases. 

The barrier height for the metal /semiconductor contacts 
can also be determined from the plot of In (I) vs V. From 
these curves the value of the saturation current l can be 
obtained by extrapolating the linear portion of the curve to 
V=0. Knowing I , Richardson's constant A, the diode cross 


sectional area S, and the temperature T, the barrier height 
(d B ) can then be determined from: 

I =SAT 2 exp(-«> B /kT) • (4.2) 

The value of barrier height obtained in this way is the zero 
bias barrier height and includes the image force barrier 
lowering A<D B . The values of * B for Au/p-ZnTe contacts are 
shown in Table 4.1. However, when the diode ideality 

Table 4.1. Barrier heights for Au/p-ZnTe contacts. 

Heat Treatment 

$b (eV) 

As -deposited 


150°C, 15 min 


200°C, 15 min 


250°C, 15 min 


350°C, 15 min 



factor, n, was determined from the following equation: 

I=I [exp(qV/nkT)-l], (4.3) 

values of =100 were obtained which violated the initial 
assumption of the analysis. Therefore this reduction of I-V 
data to <X> B is erroneous and will not be continued. 

With neither thermionic nor field emission dominating 
conduction over the voltage range from 0-3 V, it is proposed 
that thermionic field emission is the dominant mechanism. 
Thermionic field emission is an intermediate conduction 
mechanism which combines both thermionic and field emission 
processes. In a case where the barrier height is too large 
for thermionic emission to take place and the depletion 

width is too deep for tunneling to occur, current can be 
transported by the thermal excitation of holes to a level at 
which the potential barrier is thin enough for tunneling to 
occur. This process is known as thermionic field emission 
and it dominates for mid-range temperatures and moderate 
doping concentrations. In Chapter 2, the tunneling 
parameter E 00 was introduced as a measure of the relative 
importance of thermionic versus field emission transport. 
From Equation 2.13 for Au/p-ZnTe, E 00 is 0.0135 eV. This 
value is approximately the room temperature value of kT 
(0.0258 eV) which indicates thermionic field emission should 
dominate conduction. While this value of E 00 is smaller 
than kT, increasing the near-surface carrier concentration 
by Au diffusion would cause the value of E 00 to approach kT. 
The non-linearity of the ln(I) vs V plots at V<3V would 
discount pure field emission. 

TFE dominance is also supported by the fact that Au is 
a relatively deep acceptor in ZnTe (272 meV) . This level is 
too deep for the acceptors to be sufficiently activated at 
room temperature, leading to the non-linear I-V curve for 
as-deposited samples. Upon heat treatment of the Au/p-ZnTe 
contacts, increased Au diffused into the p-ZnTe would in 
turn increase the carrier concentration. The Au atoms then 
can participate in conductivity through the process known as 
hopping conduction. [Mil73] In this process, a charge 
carrier, in this case a hole, can move from one deep defect 



center to another by the tunnel effect without activation 
into the valence band. This method of conductivity is 
common where there is a high density of these centers with 
no alternative transport mechanism [Mark95] . 

Based on the AES and SIMS depth profile data, no 
compounds formed at T^250°C which could have affected 
current transport, which lends support to the Au doping 
postulate. Surface conductivity played no role in the I-V 
data measured for the contacts heat treated for 15 minutes. 
The surface diffusion, evident in Figure 4.11, occurred upon 
heating for extended times (90 minutes) and thus had an 
insignificant role in the I-V data measured after 15 minutes 
when no visible surface diffusion was present. 

The final change in the I-V curves for Au/p-ZnTe upon 
heat treating was the dramatic increase in resistance upon 
heating to 350°C. As was shown in the I-V curves of Figure 
4.3, the current passed following a 350°C heat treatment was 
approximately two orders of magnitude lower than for the 
contacts heat treated at lower temperatures. It was also 
shown that increased heat treatment time at this elevated 
temperature caused an increased resistance. It is proposed 
that widespread diffusion of Au along with interfacial phase 
formation, as evidenced from AES and SIMS depth profiles, 
limited the temperature stability of these contacts. 

Au and Te form a congruently melting compound AuTe 2 
[Mas90] which is postulated to be responsible for the 


contact degradation in these Au/p-ZnTe samples. It is 
suggested that this AuTe 2 phase is forming at the expense of 
the ZnTe layer, causing increased resistance. As mentioned 
previously, Au may be substitutionally incorporated into the 
ZnTe lattice on a Zn site. With increased time at this 
elevated temperature, more Au can be incorporated into the 
ZnTe until a point at which AuTe 2 forms in place of the 
ZnTe. It is this total breakdown of the ZnTe, caused by Au 
diffusion and formation of AuTe 2 , which is postulated to be 
responsible for the degraded contact performance at 350 °C. 
The I-V curves in Figure 4.2 show that as the heat treatment 
time at 350°C was increased, more degradation occurred which 
supports this postulate since diffusion is both time and 
temperature dependent. 


Formation of ohmic contacts to nitrogen doped p-ZnTe by 
sputter deposition of Au films was demonstrated. As- 
deposited contacts were rectifying with a breakdown voltage 
of =0.5 eV, close to the barrier height value of 0.66 eV 
obtained from the band alignment. This strongly suggests 
that the Fermi level was not pinned in ZnTe. Slightly 
decreased breakdown voltages after heat treatments at 150°C 
were attributed to the breakup of interfacial contamination 
layers. Heating for 15 minutes at T*200°C resulted in ohmic 


behavior which was credited to creation of near-surface 
acceptors due to doping by Au without interfacial phase 
formation. This led to increased conductivity through the 
near -surface layer, culminating in a maximum current density 
of 2.3 A/cm 2 at 5 V following 15 minutes at 250°C. This 
increased conduction was attributed to thermionic field 
emission occurring in the contacts at V<3V, and field 
emission for V>3V. Au acceptors with an energy of Ev+272 
meV were proposed to participate in current transport by 
hopping conductivity between deep centers in the lattice. 
Diffusion of Au was extensive and resulted in AuTe 2 compound 
formation at 350°C, leading to severe degradation of the 
contact properties. 



1.0e-2 - 

rr 5.0e-3 - 





0.0e+0 - 

-5.0e-3 - 

-1.0e-2 - 


Figure 4.1. I-V curves for Au/p-ZnTe. 






4.0e-4 - 

r? 2.0e-4 - 

O.Oe+0 - 

-2.0e-4 - 

-4.0e-4 - 







■ ' 














Voltage (V) 

Figure 4.2. I-V curves for Au/p-ZnTe heat treated at 
350°C in forming gas. 



1.0e-2 - 

— 5.0e-3 - 



O.Oe+0 - 

-5.0e-3 - 

-1.0e-2 - 


-6 -4 

m um ili um 111 IITn 

• As-deposited 
° j 250°C, 15 min 
■ : 250°C, 90 min 
D 350°C, 15 min 

i i r 

■2 2 

Voltage (V) 

Figure 4.3. I-V curves for Au/p-ZnTe, 



2.5e+4 - 

>, 2.0e+4 



<D 1.5e+4 - 



1.0e+4 - 

5.0e+3 - 


r*-"^*^ -* fca^a A 


fcftAAAjaa -* 


7 Se 

~^-^p«y a ' ■- — -, | — M 


10 15 20 25 30 35 40 

Sputter Time (Min) 

45 50 

Figure 4.4. AES depth profile of Au/p-ZnTe, as-deposited. 



2.5e+4 - 


2.0e+4 - 






1.5e+4 - 

1.0e+4 - 

5.0e+3 - 


0.0e + lA k*^..*s -l.j^w*-,a — a-^a »_■ j^v„ 

10 15 

Sputter Time (Min) 


Figure 4.5. AES depth profile for Au/p-ZnTe sample heat 
treated at 200°C for 90 minutes. 


2.5e+4 - .v--'"v-..-.:\ ;V 

2.0e+4 - 

•h 1.5e+4 - 





1.0e+4 - 

5 . Oe+3 - 

0. Oe+0 




-a,#i. , .a, ^ i/- A...ai^ .. .wd/j 
15 20 25 

Sputter Time (Min) 

Figure 4.6. AES depth profile for Au/p-ZnTe sample heat 
treated at 250°C for 90 minutes. 



3e+4 - 



C 2e+4 

le+4 - 

Oe+0 - 




^ ^- -V- ./3r-.A-^ 



Sputter Time (Min) 

Figure 4.7. AES depth profile for Au/p-ZnTe sample heat 
treated at 350°C for 90 minutes. 


Sputter Time (Min) 

Figure 4.8. SIMS depth profile for Au/p-ZnTe sample, 



le+5 - 

h le+2 - 


ui^ .y./ 

i e +o - n « J I'J 

VjA ZnTe 

~i 1 

5 10 

Sputter Time (Min) 



Figure 4.9. SIMS depth profile for Au/p-ZnTe sample heat 
treated at 250°C for 90 minutes. 




le+5 - 


M n rfM 


1 I 

5 10 15 

Sputter Time (Min) 

1 Oi It i < ; 


Figure 4.10. SIMS depth profile for Au/p-ZnTe sample heat 
treated at 350°C for 90 minutes. 




Figure 4.11. SEM micrograph of Au/p-ZnTe sample heat 
treated at 250°C for 90 minutes showing an extended reaction 
zone surrounding the Au contact. 


Figure 4.12. SEM micrograph of the surface of a Au dot on a 
Au/p-ZnTe sample heat treated at 350°C for 90 minutes 
showing roughening of the Au contact. 






-4 - 

-6 - 

-8 - 


-12 - 




lSCC, 15 min 
200°C, 15 min 
250°C / 15 min 
350°C, 15 min 


Voltage (V) 

Figure 4.13. Plot of Ln(I) vs. V for Au/p-ZnTe, 



The formation and characterization of electrical 
contacts to p-GaN are described below. The contact 
metallizations consisted of a single element (Au) , bi-layer 
contacts (Pd/Au, Ni/Au, and Cr/Au), and a three layer 
contact scheme (Ni/C/Au) . The elements are listed in the 
order in which they were deposited, i.e. the first element 
was deposited on GaN. As for Au/p-ZnTe contacts discussed 
in Chapter 4, the effects of elevated temperature treatments 
for various times on the electrical properties of the p-GaN 
contacts were determined using I-V measurements. Analytical 
characterization of the contacts included AES and SIMS depth 
profiles, along with energy-dispersive x-ray spectroscopy 
(EDS) and cross-sectional transmission electron microscopy 
(XTEM) to determine composition and interfacial reaction 
products. Photoluminescence (PL) and Hall measurements were 
also used to determine doping character of the material. 
Temperature dependent I-V data were used to determine the 
dominant conduction mechanisms for the Ni/Au, Pd/Au and 
Cr/Au contacts. In this chapter the results from I-V 



measurements, analytical characterization, and temperature 
dependent I-V are given for all metal contacts, followed by 
a discussion of these results. Finally, the results are 



I-V results 

Au contacts (2000A) were sputter deposited onto p-GaN 
which had a resistivity of 7 Qcm and a free hole 
concentration N A =lxl0 17 cm" 3 . As deposited the I-V curves 
were rectifying with a breakdown voltage of =2.5 V (Figure 
5.1) . The barrier height (® B ) was calculated using the 
method described in chapter 4 to be 0.62 eV, but again the 
ideality factor was =100, therefore no further <t> B data will 
be reported. Instead only the breakdown voltage will be 
determined using the procedure described earlier. When this 
same sample was heat treated at 400 °C for 5 minutes in 
forming gas, the breakdown voltage was reduced to =0.8 V. 
After 600°C for 15 minutes, the breakdown voltage decreased 
only to =1.5 V. The I-V curves remained rectifying for all 
heat treatments. 

The Au and Ga signals in AES depth profiles from 
contacts as-deposited or heat treated at 400°C for 15 


minutes in forming gas can be seen in Figure 5.2. For the 
as-deposited sample, there is a very sharp interface between 
the Au and GaN with no indication of GaN decomposition. At 
400°C for 15 minutes, the slope of the Au signal is 
identical to that of the as-deposited profile. The 
interface is still planar with no evidence of GaN 
decomposition or Au diffusion into the GaN matrix to the 
level detectable by AES. 


I-V results 

Contacts of Pd/Au (500/1000A) on p-GaN (N A =9.8xl0 16 cm" 3 ) 
were also studied. For as-deposited contacts, the I-V 
curves were rectifying (Figure 5.3) with a breakdown voltage 
of =3 V. Heating at 200°C for 5 minutes in flowing N 2 
caused the contacts to became more rectifying and pass less 
current with an increased breakdown voltage. The I-V curves 
were unchanged by increasing the time at 200°C up to 30 
minutes. After 400°C for 5 minutes, the current was 
slightly greater than for as-deposited samples. Heating for 
times up to 45 minutes at 400°C led to a 60% increase in 
current and a reduction in breakdown voltage from 2 to 1.5 
V. The I-V data from samples annealed at 600 °C for 5 
minutes were similar to those after 400°C for 45 minutes. 
After an RTA at 900°C for 15 seconds in N 2 , the contacts 


still rectifying, but the breakdown voltage was decreased to 
=1 V with 5 mA of current at 5 V. 


As-deposited (Figure 5.4), the Au/Pd interface was 
broadened and at the Pd/GaN interface there was a slight Pd 
tail into the GaN, but no evidence of Ga out-diffusion into 
the overlaying films following deposition. At 200°C for 5 
minutes (Figure 5.5) there was an increased broadening of 
the Au/Pd interface with a small Pd signal continuous 
throughout the Au layer to the surface. The AES survey 
spectrum (Figure 5.6) showed a Pd transition at 327 eV, 
indicating that Pd diffused through the Au layer. The Pd 
signal also tailed into the GaN, but there was no evidence 
of Ga out-diffusion. After 400°C for 5 minutes (Figure 
5.7), Pd continued to diffuse through the Au capping layer, 
but did not build up at the surface. Thus there were 
constant levels of Pd and Au throughout the first layer, 
followed by a Pd-rich layer underneath. The Pd/GaN 
interface has broadened, but there was still no evidence of 
dissociation of the GaN. In Figure 5.8, the AES depth 
profile after an RTA of 900°C, 15 seconds shows that Pd had 
completely dissolved into the Au layer forming a Au:Pd 
region contacting the underlying GaN. The depth profile 
suggests that dissociation of GaN occurred because a Ga 
signal was detected throughout the contact layer up to the 
surface. The long Pd and Au tails into the GaN suggest 


diffusion into the GaN. Following this RTA heat treatment, 
the contacts exhibited both dark regions rich in Pd and Au 
(Figure 5.9), and light regions rich in Pd, Au, Ga, and N 
(Figure 5.10) . 


I-V results 

Evaporated Ni capped with sputtered Au (500/2400A) 
contacts deposited on the same substrate as the Au contacts 
were rectifying as-deposited with a lower breakdown voltage 
(=0.5 V) and larger current (=0.5 mA @ 5 V) than the Au/p- 
GaN contacts (Figure 5.11). The I-V curves remained 
rectifying up to 200°C, but after a 400°C, 5 minute 
treatment, they became nearly linear with =0.6 mA passed at 
5 V. Upon heating this sample to 600°C for 5 minutes, the 
contacts degraded as reflected by lower current levels 
(Figure 5.11) . 

The effects of Ni layer thickness and Ni to Au 
thickness ratios are shown in Figure 5.12. Data are shown 
from as-deposited and heat treated (400°C, 5 min. in forming 
gas) samples (p=17 Qcm) with evaporated 200/2000A or 
500/1000A Ni/Au contacts. The as-deposited samples were 
very similar in both current magnitude (0.1 mA @ 5 V) and 
breakdown voltage (=2.2. V). After heat treatment, the 
breakdown voltage was drastically reduced to =0.5 V even 
though the I-V curves remained slightly rectifying. There 



was very little difference in the I-V data for different 
metal thicknesses and thickness ratios after the heat 
treatment . 

The effects of heat treatment ambient were also 
investigated for Ni/Au contacts. Initial heat treatments 
were performed in a forming gas environment, but this was 
changed to N 2 due to reports of hydrogen passivating Mg 
acceptors in p-GaN [Pea96a] . Use of N 2 also reduced the 
probability of N vacancy formation in GaN which may lead to 
semi -insulating films. The difference in current levels 
obtained was negligible between the different ambients, but 
it was determined that N 2 were preferable to prevent any 
potential acceptor passivation, especially at elevated 
temperatures . 

An AES depth profile for an (500/1000 A) contact can 
be seen in Figure 5.13. There is a planar Ni/Au interface 
with no indication of intermixing upon deposition. The 
Ni/GaN interface appears less sharp with a slight Ni tail 
into the GaN indicative of near-surface diffusion, but this 
could result from loss of resolution at greater sputter 
depths [Hol81] . There is, therefore, no evidence of GaN 
dissociation upon deposition except for the rather low 
breakdown voltage for as-deposited films. The depth profile 
for a sample heated to 200°C for 5 minutes in flowing N 2 was 
similar to the as-deposited profile. After 400°C for 5 


minutes (Figure 5.14) the Ni/GaN interface was slightly more 
diffuse, and Ni diffused through the Au capping layer to the 
surface of the contact. For heat treatments of 600°C 
(Figure 5.15), dramatic evidence of GaN dissociation is 
obvious. Ga was detected at the Au/ambient surface, however 
Ni also diffused through the Au capping layer to the surface 
and formed a NiO x layer. There was also an increased C 
signal at the metal /p-GaN interface. 

Ni / C/Au 

I-V results 

As will be discussed, the increased current for heat 
treated Ni/Au contacts could have resulted from C doping of 
the GaN surface. Therefore, the next contact scheme 
investigated was an evaporated carbon layer intermediate 
between the Ni and Au layers to give a Ni/C/Au contact 
(500/100/1000A) on p-GaN with a carrier concentration of 
2xl0 17 cm" 3 . I-V data for the Ni/C/Au contacts were 
rectifying as-deposited and remained rectifying with all 
heat treatments up to 400°C for 5 minutes (Figure 5.16) . 
The contacts exhibited a decrease in current vs. voltage 
upon heating to 200°C for 5 minutes. This was then followed 
by increased current at 400°C for 5 minutes. For all heat 
treatment conditions, the currents were lower and the 
potential offsets higher than in the Ni/Au contacts under 
similar processing conditions. 



The AES depth profile from an as-deposited Ni/C/Au 
contact (Figure 5.17) showed more diffuse Au/C, C/Ni and 
Ni/GaN interfaces than for a Au/Ni contact. The peak of the 
carbon signal occurred before the Au signal had decreased to 
zero and there was a very long C tail into the Ni layer. 
The Ni/GaN interface was also very broad. The depth profile 
after heat treatment at 400°C for 30 minutes (Figure 5.18) 
was similar to the as-deposited profile with no evidence of 
Ni diffusion to the surface or of GaN dissociation. 

In the SIMS depth profiles, there were well defined 
interfaces between all layers in an as-deposited sample 
(Figure 5.19) with a noticeable C tail into the Ni region of 
the sample. After heating to 400°C for 30 minutes in 
forming gas (Figure 5.20), the interfaces were less well 
defined and the C signal at the Au/Ni interface broadened. 
As in the Ni/Au contacts, there was evidence of Ni diffusion 
through the Au contact to the surface and the Ni signal 
tailed into the p-GaN substrate with no evidence of Ga at 
the ambient/Au surface. 


I-V results 

I-V results from as-deposited Cr/Au (500/1000A) 
contacts were once again rectifying, but with a breakdown 
voltage =1-1.5 V (Figure 5.21). Annealing at 200°C for 5 


minutes in flowing N 2 resulted in rectifying contacts with 
lower transported current, similar to the other contact 
schemes. At 400°C for 5 minutes, the breakdown voltage was 
-IV. Following an RTA at 900 °C for 15 seconds, the I-V 
curves were linear with =13 mA @ 5 V. For this sample an 
upper limit on the specific contact resistance has been 
calculated to be i4.3xl0" 1 Qcm 2 by assuming a bulk resistance 
of zero for the underlying p-GaN. 

In the as-deposited AES depth profile (Figure 5.22), 
the Cr/Au interface was sharp while the Cr/GaN interface 
showed a Cr tail into the GaN, but no evidence of 
dissociation of GaN. In contrast to the AES data from the 
other metallization schemes, there were small but detectable 
carbon concentrations at the Cr/Au and Cr/GaN interfaces. 
After 200°C for 5 minutes in N 2 (Figure 5.23), the depth 
profile was similar in all aspects to the as-deposited 
profile. At 400°C for 5 minutes, Cr diffused through the Au 
capping layer to the surface and formed an oxide layer 
similar to the Ni reaction observed in the Ni/Au contacts 
(Figure 5.24 vs Figure 5.14). Following this heat 
treatment, an increased C signal was only detected at the 
Cr/GaN interface. After a 900°C, 15 second RTA (Figure 
5.25), there was massive Cr diffusion through the Au capping 
layer to the surface and also into the GaN layer. GaN 
dissociation was evident from the large Ga signal in the Au 


region (-8-15 minutes sputter time) without a corresponding 
N signal and the small Ga signal at the surface. There were 
increased Cr, nearly constant N and reduced Ga signals in 
the region of the original Cr layer. 

XTEM and EDS were performed on as -deposited Cr/Au 
contacts and following the 900°C, 15 sec RTA anneal. For 
the as-deposited sample, the micrograph in Figure 5.26 shows 
a low magnification (x 73K) image of the contact/GaN 
structure. The polycrystalline Au and Cr metal layers are 
fairly uniform in thickness across the contact. In this 
figure the appearance of the Au layer being <1000A thick is 
attributed to the ion milling process which removed part of 
the Au capping layer. The ring structure of the selected 
area diffraction pattern (SADP) in Figure 5.27 is indicative 
of polycrystalline material with rings from both Au and Cr 
observed as shown. EDS data (Figure 5.28) were consistent 
with discrete metal layers. After the RTA anneal, the 
metal /GaN interface was no longer planar as shown by Figure 
5.29. In the thicker regions of the metal contact in Figure 
5.29, EDS spectra (Figure 5.30) show the presence of both Au 
and Cr. While exact compositions can not be obtained from 
this data, the relative intensities of the Au and Cr peaks 
can be used to ascertain that Cr has diffused through the Au 
capping layer and there are regions where both Au and Cr are 
present which is consistent with intermixing. This Cr 


diffusion is also supported by the depth profile (Figure 
5.25) . 

In Figure 5.31 a high resolution XTEM of the 
ambient/contact surface is shown. The presence of Ga at 
this surface, as determined by EDS (Figure 5.32), indicated 
Ga out-diffusion from the GaN and segregation to the 
surface. SADP's taken from the regions indicated in Figure 
5.31 show a polycrystalline phase containing Au and Cr 
(Figure 5.33), and a multi-crystal phase which includes Ga 
along with Au and Cr (Figure 5.34). 

Temperature Depe ndent I-V 

Temperature dependent I-V measurements were used to 
determine the conduction mechanisms dominating in the Ni/Au, 
Pd/Au and Cr/Au contacts. The I-V data for a Ni/Au 
(500/2500A) sample heated to 600°C (873K) for 5 minutes in 
flowing N 2 are shown in Figure 5.35. The contacts were 
rectifying at all temperatures between 80-400 K. The 
current increased and the reverse bias breakdown voltage 
decreased with increasing temperature. 

The temperature dependent I-V data for a Pd/Au 
(500/1000A) sample heated to 600°C (873K) for 15 minutes in 
flowing N 2 are shown in Figure 5.36 for temperatures 90-390 
K. As for Ni/Au contacts, the current increased and the 
reverse bias breakdown voltages decreased with increasing 
temperature. The current increase was much more dramatic 


for Pd/Au versus Ni/Au contacts, but all contacts remained 

For Cr/Au (500/2500A) contacts annealed at 400°C (673K) 
for 5 minutes in flowing N 2 , the 80-400 K temperature 
dependent I-V data can be seen in Figure 5.37. The I-V 
curves were slightly rectifying at all measurement 
temperatures. Unlike Pd/Au and Ni/Au contacts, there was 
very little increase in current with increased temperature. 
The reverse bias breakdown voltage also remained nearly 
steady over this temperature range. 

Disc us sion 

Metal Contact Schemes 

As previously mentioned in Chapter 1, GaN has a direct 
band gap of 3.4 eV and an electron affinity between 3.3-4.1 
eV. Plugging these values into Equation 2.1 to find the 
potential offset in the absence of Fermi level pinning leads 
to values of 6.7-7.5 eV required for metal work functions to 
decrease the potential offset at the semiconductor valence 
band to zero. This becomes important due to the fact that 
metal work functions do not have values greater than =5.5 eV 
[Sze81] . Thus in the absence of interfacial reactions, 
there will always be a potential offset between metals and 
p-GaN. Because of this, the interfacial reactions between 
the metal contacts and p-GaN upon heat treatment are of 


utmost importance for formation of low resistance ohmic 

Thus there was a need to discover a metal contact 
scheme that would react with the underlying GaN either as- 
deposited or upon heat treatment to form ohmic contacts. 
These metal schemes could provide low resistance contacts in 
one of two ways. First, the formation of a large work 
function compound or phase would decrease the potential 
offset at the valence band, and second the doping 
concentration in the near-surface region of the 
semiconductor could be increased by incorporation of 
acceptor dopant species from either the contact layer or 
interfacial contamination. 


The work function of Au is 5.1 eV [Sze81], so the 
calculation reported above predicted that the as-deposited 
contacts would be rectifying, as was the case. Au has been 
reported to exhibit limited reaction with GaAs as a single 
element contact and increased reactivity upon incorporation 
in a contact scheme such as Au/Ge/Ni or Au/Zn [Bras83] . The 
ohmic behavior is then observed through a solid state 
regrowth process in the manner described by Sands et al . 
[San88], and Holloway et al. [Hol97] . As seen in Figure 5.2 
for the Au/GaN contacts, there was no evidence of Au:Ga 
phase formation upon deposition or following heating at 
400 C C. Thus Au on GaN is non-reactive, leading to the 


conclusion that the presence of additional metals in the 
contact scheme is required for ohmic contact formation. 


With no discernible Au/GaN reaction evident upon heat 
treatment, inter facial metal layers were introduced into the 
contact scheme to induce reaction with GaN. For Pd/Au 
contacts (500/lOOOA) interfacial reactions played extremely 
important, complex roles in interface transport. As 
reported above, there was evidence of a small uniform amount 
of Pd throughout the Au contact layer after heating to 
200°C, but surface segregation was not observed. Upon 
heating at 400°C for 5 minutes, diffusion continued until a 
signal intensity ratio of -2:1 Au:Pd formed at the surface, 
as can be seen in Figure 5.7. The I-V data showed higher 
current levels and lower breakdown voltages at 400°C. It is 
proposed that as a result of the heat treatment there was 
interface dissociation and breakup of the interfacial oxide 
on GaN which led to increased conduction across the 
contacts. This is consistent with reports that Pd will 
dramatically dissolve adsorbed C and 0. This reaction 
continued with increased heat treatment time at 400 °C, 
leading to lower contact resistances. This reaction is 
presumed to be completed at 600°C since longer times at this 
temperature did not lead to increased conduction. Upon an 
RTA heating to 900°C for 15 sec, the current levels 
increased to their greatest values and this increase was 


accompanied by dissociation of the GaN as evidenced by the 
increased Ga signal throughout the entire Au:Pd contact 

Thus, the breakdown of the inter facial oxide and 
contamination caused by the dissolving of C and in the Pd 
layer is believed to be responsible for the increased 
conduction upon annealing at temperatures between 400° and 
600°C. It is speculated that upon heating to 900°C, 
dissociation of GaN occurred allowing increased doping of 
the near-surface region, possibly by Pd acceptors or 
inter facial C contamination. At lower temperatures, with no 
GaN dissociation, a lower Ga vacancy concentration decreased 
the probability for acceptor species to incorporate 
substitutional^ in the GaN lattice. Thus while Pd aids in 
the dissociation of the interfacial oxide, it may lessen the 
effects of C doping of the near-surface region and does not 
lead to an ohmic contact. 

The next contact scheme investigated included an 
interfacial Ni layer to induce reaction with GaN leading to 
increased doping. This is the contact scheme presently used 
in experimental laser diodes [Nak97] . Bermudez et al . have 
deposited Ni layers on atomically clean GaN and investigated 
the reactions [Ber93] . They determined that chemical 
reaction at the interface between GaN and Ni occurred even 
near room temperature, and annealing above =600 °C led to a 


pronounced inter facial reaction with the appearance of Ga at 
the Ni surface. They also determined that the Ga remained 
dissolved in the Ni layer rather than forming a stable 
compound or surface segregated layer. 

In this study, Ni (<t> Ni =5.15 eV) formed a rectifying 
contact as-deposited with a breakdown voltage of 0.5-0.6 eV. 
This value is lower than that predicted from the energy band 
alignment (1.55 eV) for a clean GaN surface with no Fermi 
level pinning which would be consistent with undetected 
dissociation of GaN upon deposition. In the AES depth 
profile from an as-deposited contact (Figure 5.13), the 
slight Ni tail into the GaN is consistent with Ni diffusion 
into the semiconductor during deposition, as suggested by 
Bermudez et al. [Ber93] . With heat treatment to 400°C, 
there was evidence of Ni diffusion through the Au capping 
layer to form NiO x at the surface. 

At 600°C for 30 minutes, the AES depth profile (Figure 
5.7) showed direct evidence of GaN dissociation upon heat 
treatment. Dissociation of the GaN was evident from the 
fact that Ga diffused through the Ni/Au contact layer and 
piled up at the surface. The Ni and Au interdif fused and 
the width of the Ni/GaN interface increased significantly, 
consistent with widespread diffusion of Ni into the GaN 
lattice or roughening of the Ni/GaN interface. The reaction 
products of dissociation of GaN by Ni are uncertain, since 
there is a low solubility of N in Ni even though Ni will 


combine with Ga to form a solid solution and several binary 
intermetallic phases [Mas90] . Unfortunately, simple 
dissociation of the GaN lattice should not result in ohmic 
contacts without an increased carrier concentration in the 
near surface GaN region. There are two potential acceptors 
in this system since Ni has been shown to be a deep acceptor 
in GaAs [Mil73], thus the possibility exists that Ni will 
behave similarly in GaN. Carbon from interfacial 
contamination may have incorporated into the GaN lattice 
during this interfacial reaction, and Abernathy et al . 
[Abe95] have shown that C can act as an acceptor in GaN. 
Carbon was found to form a shallow acceptor with a binding 
energy of -0.2 eV while sitting on the nitrogen sublattice 
[Fis95] , therefore it is reasonable to speculate that 
interfacial contamination detected by both AES and SIMS 
could lead to a higher free hole concentration at the 
interface. It has been calculated that in order to obtain 
field emission in these Ni/Au on p-GaN contacts, a doping 
density of ~2xl0 19 cm" 3 would be required. Thus the 
combination of interfacial C and possibly Ni acceptors could 
provide the needed carriers for ohmic contacts to form. In 
summary, diffusion and intermixing of Au and Ni occurred 
upon heating to 400°C, with extensive dissociation of GaN at 
600°C, and thus led to improved contact properties. 



Contacts of Ni/C/Au, with 100A of evaporated C were 
tested to determine if increased doping and better ohmic 
contacts would be observed with the intentional presence of 
C. As shown in Figure 5.8, this contact scheme exhibited a 
higher resistance than did the Ni/Au scheme under all heat 
treatment conditions. Additional evaporated C at the 
metal /semiconductor interface did not improve the contact 
behavior. This could result from a number of factors, 
including that the concentration of interfacial C 
contamination was sufficient to saturate the acceptor doping 
effects and the additional evaporated C layer led to self- 
compensation of carbon acceptors (C N ) by carbon donors (C Ga ) . 
In addition, interfacial contaminating carbon is largely 
"adventitious" organic and hydrocarbon molecules which are 
similar to the C source used by Abernathy et al. [Abe95] to 
demonstrate p-doping of GaN. Evaporated C in a thin film is 
bound primarily in a ring structure, which is notoriously 
poor as a doping source during thin film growth. 

With regard to Cr/Au (500/1000A) contacts, Cr has a 
smaller work function (<& Cr =4.5 eV) than the other metals 
investigated, so there is a greater need for interface 
reactions. As with all bi -metal contacts, the two important 
interfaces are Cr/GaN and Cr/Au. As-deposited XTEM (Figure 
5.26) showed uniform metal layer thickness across the 


contact, and the AES depth profile (Figure 5.22) showed an 
abrupt Cr/Au interface. However, the Cr/GaN interface was 
broadened with an elevated C level, but no GaN dissociation 
could be detected. This broadening of the Cr/GaN interface 
is most likely due to two factors, the diffusion of Cr into 
GaN up to the solubility limit, and the roughening of the 
interface that occurs due to differential sputtering of the 
layers. For furnace anneals up to 400°C, there was no 
change in the Cr/GaN interface when compared to the as- 
deposited case. This was in contrast to the Cr/Au interface 
which did not noticeably broaden to levels detectable in AES 
depth profiles at 200°C, but at 400°C Cr diffused through 
the Au capping layer to form a surface oxide layer. This is 
similar to the case of Ni in the Ni/Au contacts. The I-V 
data for all these samples were similar to the as-deposited 
data in current magnitude. Overall, the diffusion of Cr to 
the Au/ ambient surface did not result in better contact 
properties, presumably due to the lack of continued reaction 
at the Cr/GaN interface at the temperatures studied. 

While the Cr/Au contact results were rectifying for 
anneal temperatures of 400°C or less, a 900°C, 15 second RTA 
resulted in ohmic contacts (Figure 5.21). The specific 
contact resistance reported above was calculated from the 
current (11.4 mA) at 5V and using Ohm's law, with the 
resistance being multiplied by the area of the dot contacts 
(1.96xl0~ 3 cm 2 ) to give p c =4.3xl0" 1 Qcm 2 . In this procedure, 


the bulk resistance of the GaN is assumed to be zero. Since 
p GaN is finite, the actual specific contact resistance is 
lower than the value reported above. In any case, p c <10" 4 is 
desired for practical devices. 

For the RTA sample, an AES depth profile (Figure 5.25) 
showed widespread diffusion of Cr throughout the entire 
structure along with GaN dissociation. The varying layer 
thickness across the contact evident in the XTEM micrograph 
(Figure 5.29) is due to the intermixing of Au and Cr which 
is also shown in the depth profile. As opposed to the as- 
deposited sample, ion milling did not effect the thickness 
of the metal layers as shown by the remaining glue at the 
metal surface. In this contact scheme, much like the Ni/Au 
scheme, dissociation of GaN took place with Ga diffusion 
through the contact layers. The Ga diffused to the surface 
in some regions as evidenced by EDS data, and Ga also 
segregated into the Au region of the contact as was seen in 
the AES depth profile. 

For the RTA sample, both a AES depth profile (Figure 
5.25) and a XTEM micrograph (Figure 5.29) coupled with EDS 
data (Figure 5.32) indicated Ga levels at the surface of the 
contact without comparable levels of N, which is consistent 
with GaN dissociation. There was however, one discrepancy 
between the two sets of data. While the XTEM and EDS showed 
a strong Ga signal at the surface with no evidence of Ga 
throughout the contact layer, the AES data indicated a 

chromium oxide with only a slight Ga signal at the surface, 
and a large accumulation of Ga in the original Au region of 
the contact. This variation can be explained in terms of 
spot size and sampling area of the two techniques. In the 
EDS, the spot size was -500A which allowed for the effects 
of surface roughness to be negligible for this measurement. 
On the other hand, the spot size for the AES was -lum. Over 
this large an area, the surface roughness evident from 
Figure 5.29 leads to data collection from different depths 
relative to the original surface and the formation of 
profiling artifacts which complicate data analysis. 

To understand the mechanism responsible for the 
increased conduction, it is important to first look at the 
properties of Cr in GaAs and relate them to Cr in GaN. In 
GaAs, chromium is a very deep acceptor (Ev+0.79 eV) that is 
generally used to make semi-insulating GaAs [Mil73] by 
substituting on the Ga site in the lattice [Bro78 and Eu80] . 
Yu et al. determined that diffusion of Cr atoms in GaAs 
results from a substitutional Cr hopping into an 
interstitial site, rapid migration of the interstitial atom, 
and a subsequent return to a vacant Ga site [Yu91] , thus the 
Ga vacancy concentration and vacancy generation rates are 
key factors in this substitutional-interstitial conversion 
process [Dea86] . Based on these data from GaAs, it is 
proposed that the 900°C, 15 second RTA caused diffusion of 
Cr into GaN. Since Cr acts as an acceptor when substituted 

on the Ga sublattice in GaAs, it is reasonable to believe it 
would behave similarly in the III-V material GaN. This is 
consistent with the fact that Cr diffused into the GaN at 
400°C which led to lower breakdown voltages (-1V) than 
predicted from the Schottky model (2.2-3.0 V). At 900°C, 
GaN dissociation occurred with out-diffusion of Ga resulting 
in a Ga vacancy flux which when combined with interstitial 
Cr diffusion, could create substitutional Cr deep acceptors 
over the depletion depth and increase the near-surface 
hopping conduction. Contrary to the GaAs/Ni/Ge/Au system 
[KimT90], there was no evidence of solid phase regrowth of 
the GaN at the contact interface in any of the XTEM 
micrographs. Thus the increased doping levels are 
attributed to Cr interstitial in-diffusion as opposed to 
solid phase regrowth with interstitial transport enabling 
this mechanism. 

In these contacts, it has been proposed that 
substitutional acceptor Cr impurities in the GaN are 
contributing to the conductivity by increasing the carrier 
concentration in this region. Presumably, the Cr impurity 
level in GaN will be quite deep and thus the problem arises 
as to how the Cr acceptors will participate in conduction 
since at room or device operation temperature there will not 
be enough thermal energy to excite a large percentage of 
carriers between the Cr acceptor level and the valence band. 
The method of conduction proposed for this situation in 

hopping conduction between Cr acceptor levels without 
excitation into the valence band. 

For hopping conduction to occur, there is an activation 
energy, usually smaller than the band to band activation 
energy, that must be overcome. This energy arises due to 
the dispersion of impurity levels. By increasing the 
impurity concentration, this activation energy can be 
decreased until eventually it vanishes which corresponds to 
the transition between activated and metallic conduction. 
The concentrations that were dealt with in this study were 
far below this level. The physical reason for the 
exponential dependence of carrier hopping results from the 
probability of a jump between two impurities being 
determined by the wave function overlap. The wave function 
overlap integrals drop exponentially with increasing 
distance between impurities. As the concentration 
decreases, the mean inter-impurity separation increases, 
with the jump probability and hence the electrical 
conductivity decreasing exponentially [Shk84] . Hopping 
conductivity has been shown in a variety of wide band gap 
semiconductors including ZnSe [Mark95] and GaAs [Loo90] . 

Another potential explanation, which will be shown to 
be inaccurate, for the observed increase in conduction in 
the contacts following the 900 C C heat treatment was that 
elevated temperature caused the formation of N vacancies in 
the GaN which led to a type shifting from p- to n-GaN. If 

this were true, a low resistance contact to n-GaN was formed 
as opposed to the desired contact to p-GaN. A recent report 
has shown that through the use of an AlN capping layer, 
ohmic behavior was reduced in n-GaN as opposed to an 
uncapped sample upon heating to 1100°C. This was attributed 
to the formation of N vacancies in the uncapped sample which 
increased the near-surface donor concentration [Zol96b] . 

To determine the effect of heat treatment on the doping 
character of the GaN, photoluminescence (PL) and Hall effect 
measurements were performed on a sample as -grown, and 
following a 900°C, 15 second RTA in flowing N 2 . The PL 
spectra from 350-650 nm (3.54-1.90 eV) for these samples can 
be seen in Figure 5.38. In this figure, the peak of the 
strongest transition is at 2.83 eV (438 nm) for the as-grown 
sample and 2.93 eV (423 nm) for the heat treated sample. 
This emission is associated with a deep Mg complex which has 
been previously determined to have an energy of -2.87 eV 
(432 nm) [Myo96] . This data shows the continued presence of 
these Mg complexes in the material . Another important 
region of the spectrum is near the band gap of GaN. For 
this region, Figure 5.39 shows the photo emission from 355- 
380 nm (3.49-3.26 eV) . The peak in each spectrum at 3.44 eV 
(360 nm) is attributed to the band edge excitation from the 
valence band to the conduction band, while the peak at 3.33 
eV (372 nm) is attributed to an acceptor bound excitation 
emitted from a Mg acceptor site. More importantly, the two 

curves are very similar in this low wavelength region and 
there is no large donor band emission in either sample that 
would indicate n-type conductivity in the material. The 
donor Si in GaN has a binding energy of ~30 meV which should 
result in a peak at ~363 nm; none is observed either before 
or after RTA. 

To further demonstrate that the 900°C, 15 second RTA 
did not convert the surface to n-GaN, Hall data were 
collected from a sample as-grown and after RTA. Assuming a 
doped layer thickness of 1000A, the calculated free carrier 
concentration was 2.6xl0 18 cm" 3 as-grown and 4.5xl0 18 cm" 3 
following the heat treatment, in both cases p-type. This 
showed that the material did not become n-type following the 
heat treatment and actually had more free carriers than in 
the as-grown case. This can be attributed to either 
experimental error or some activation of Mg acceptors at 
this temperature, possibly by dissociating H-Mg complexes. 
Thus it has been shown that any conversion from p- to n-GaN 
would require a higher heat treatment temperature and 
possibly a different heat treatment environment. 

Temperature Depe ndent I-V 

Temperature dependent I-V data were used to determine 
the dominant transport mechanism (s) in the Ni/Au, Pd/Au and 
Cr/Au contacts to p-GaN. As described in Chapter 2, the 
tunneling parameter E 00 can be used as a measure of the 


relative importance of thermionic versus field emission 
transport. Since all contact metallizations were deposited 
on separate pieces of GaN grown on the same piece of 
sapphire, the theoretical value of E 00 should be constant 
for each of the metallizations. From Equation 2.13, the 
theoretical value of E 00 is 0.00223 eV for contacts to GaN. 
For ohmic contacts it is preferable that E 00 »kT. 
Unfortunately, for these contacts, E 00 «kT for all 
temperatures investigated, which results in the prediction 
that thermionic emission should dominate conduction in each 
of the contact schemes. This is understandable since the 
hole concentration of the GaN is relatively low (9.8xl0 16 
cm" 3 ) . However, if the near-surface carrier concentration 
was increased due to inter facial reactions, the dominant 
conduction mechanisms would be expected to change. 

For contacts dominated by thermionic emission of holes 
over the valence band offset, the saturation current is 
described by Equation 2.7 which can be rearranged in the 
following way: 

ln(I /T 2 )=ln(SA*) + (A* B -<D B )/kT. (5.1) 
From this equation, plotting the temperature dependent I-V 
data as ln(I /T 2 ) vs 1/T should produce a straight line if 
thermionic emission was the only mechanism resulting in 
conduction. These data are plotted for select voltages for 
each of the contact schemes from 100-400K in Figure 5.40 and 
from 300-400K in Figure 5.41. On the other hand, if 

conduction is dominated by field emission, Equation 4.1 can 
be used where a plot of ln(I) vs V will be linear; this is 
shown from 0-5 V in Figure 5.42, and from 3-5 V in Figure 

For the Ni/Au contacts Figure 5.35 showed that at a 
constant voltage, current increased with increasing 
temperature, this is consistent with thermionic emission. 
From Figures 5.40-5.43, the data indicates that thermionic 
emission is dominating conduction for low voltages and 
normal operating temperatures (300-400 K) , whereas at 
reverse bias 2:3.0 V there is an increased field emission 
component . 

For the Pd/Au contacts it is important to note that the 
sample analyzed was heat treated at 600 °C, a temperature at 
which the GaN was not expected to dissociate. This was also 
not the optimum heat treatment in terms of limiting 
breakdown voltage. Figure 5.36 indicates thermionic 
emission dominates due to the very low current levels at 90 
K and the tremendous increase at 390 K. From the data in 
Figures 5.39-5.42, it is evident that at low values of 
reverse bias thermionic emission is dominating the 
conduction. However, upon reaching a value of =3 V, field 
emission currents play a role in the conduction. This 
impact is smaller than in the Ni/Au contacts. 

The Cr/Au sample was also heat treated at 600°C, a 
temperature at which a non-linear room temperature I-V curve 

was obtained. However, this contact scheme showed that 
tunneling dominated more than for the other schemes. Figure 
5.37 showed nearly temperature independent I-V curves over 
the temperature range 80-380 K, which suggests control by 
field emission. From Figure 5.40, the ln(I /T 2 ) vs 1/T plot 
showed slight but steady curvature at a constant temperature 
as the voltage was increased, and in Figure 5.42 for Cr/Au, 
linear In (I) vs V was obtained for 2-5 V. It is thus 
believed that field emission currents play a much larger 
role in the Cr/Au contacts than the other contact schemes 
and for V*3V, this mechanism clearly dominates transport. 


Formation of electrical contacts to magnesium doped p- 
GaN by sputter deposition of Au and electron beam deposition 
of Pd/Au, Ni/Au, Ni/C/Au and Cr/Au thin film contacts was 
demonstrated. The best contact schemes to p-GaN took 
advantage of interfacial reactions between the metals and 
the semiconductor, in particular dissociation of GaN leading 
to Ga out-diffusion. Pure Au contacts showed no reaction 
with GaN up to 400°C and resulted only in rectifying 
contacts. Pd/Au contacts were rectifying under all process 
conditions and increased conduction was attributed to 
dissociation of the GaN following a 900°C 15 second RTA, 
with Pd or C potentially acting as an acceptor. Insertion 

of a Ni inter facial layer to form a Ni/Au contact led to 
dissociation of the GaN at 600°C and while ohmic contacts 
were not obtained, the nearly-linear I-V curves were 
attributed to increased doping of the GaN near-surface 
region by interfacial carbon contamination or possibly Ni 
acting as an acceptor. Introduction of an intermediate C 
layer between Ni and Au to form a Ni/C/Au contact did not 
improve the contact resistance. Films of Cr/Au provided the 
best ohmic contacts with p c i4.3xl0" 1 Qcm 2 following a 900°C, 
15 second rapid thermal anneal. At this temperature Cr 
caused the dissociation of GaN and was postulated to act as 
a substitutional acceptor sitting on the Ga sublattice. GaN 
doping was postulated to occur by an interstitialcy 
diffusion mechanism based on similar data for Cr in GaAs. 
Temperature dependent I-V measurements showed Ni/Au, Pd/Au 
and Cr/Au to all have components of both thermionic emission 
and field emission, with Cr/Au being dominated by field 
emission. This is consistent with the dissociation of GaN 
by Cr and enhanced diffusion to create increased doping in 
the near-surface region. 





2e-4 - 

^ le-4 - 

0e+0 - 

-le-4 - 

-2e-4 - 


* As -deposited 
°; 400°C, 5 min 
■ 600°C, 15 min 

1 1 1 r 

-6 -4-2 2 

Voltage (V) 



Figure 5.1. I-V curves for sputtered Au contacts to p- 
GaN heat treated in forming gas. 



§ 2e + 4 



le+4 - 


/ , -fN. A- A" v \-vv, N -Vv./S 1 ,^, \j\,v\ ^— s. w» 



400°C, 15 minutes 

a...a.-.*.\ ..,.wU-._^>- •■•^. i — cak^dUfci -a.'.-.~-L>.\ -A 



Sputter Time (Min) 


Figure 5.2. AES depth profiles for Au/p-GaN contacts as- 
deposited and heat treated for 15 minutes in forming gas. 





oe-j — 
4e-3 - 


2e-3 - 

Oe+0 - 
2e-3 - 

/^^ • 


y^F O 

200°C, 5 min 

4e-3 - 

^ T ™ 

400°C, 5 min 
600°C, 5 min 


900°C, 15 min 


1 1 1 




■2 2 

Voltage (V) 

Figure 5.3. I-V data for Pd/Au on p-GaN. 







1.8e+4 -r 
1.6e+4 - 
1.4e+4 - 
1.2e+4 - 
l.Oe+4 - 
8.0e+3 - 
6.0e+3 - 
4.0e+3 - 
2.0e+3 - 
0.0e+0 - 

i Au 

- ioh .1.41 



ftfwp-iy 4 

10 15 20 

Sputter Time (Min) 


Figure 5.4. AES depth profile for Pd/Au on p-GaN, 



2.0e+4 - 


■H 1.5e+4 - 


£ 1.0e+4- 

5.0e+3 - 



10 15 20 

Sputter Time (Min) 


Figure 5.5. AES depth profile for Pd/Au on p-GaN, heat 
treated at 200°C / 5 minutes in N 2 . 








H h 

H 1 1 1 1 h 

H h 

H 1- 


271 «V 

ka 89 «V 

H 1 1 1 h 


1— H 1 1— 

888 800 1868 1288 


l I 

—I 1 h 

1888 2836 

Figure 5.6. AES surface survey spectrum for Pd/Au on p-GaN 
heat treated at 200°C for 5 minutes, showing the Pd 327 eV 
peak at the surface. 


1.5e+4 - 

5.0e+3 - 







fc-«»-i*-i— *»-.*. »Y».aA.-.*» *-.>■<■. . ,t.A-^/^' 

Sputter Time (Min) 

Figure 5.7. AES depth profile for Pd/Au on p-GaN, heat 
treated at 400°C, 5 minutes in N„. 


1.2e+4 - 


Au | ;;i; 


"Ji .a> 



8.0e+3 - 


"• V 



i /' 


Xh/ N 


4.0e+3 - 


0.0e+0 - 

1 1 1 

10 15 20 

Sputter Time (Min) 


Figure 5.8. AES depth profile for Pd/Au on p-GaN, heat 
treated at 900°C, 15 seconds in N,. 





H — ^ 

- 73 Au 

-M — i — ^ 

h — i — h 

h — i — i — i — i — i — i — i — i- 



330 Pd 

H h 

H h 

H 1 H 

H 1 h 

) 1866 1296 m m 188G 2906 


Figure 5.9. AES surface survey spectrum for Pd/Au on 
p-GaN, 900°C, 15 seconds in N 2 , showing only Au and Pd 
present in this area. 





H h 

H 1 h 

H 1 1 1 


385 511 

330 Pd 

■I h 

— I 1 1 1 1 1— 

«e m m m m 


Figure 5.10. AES surface survey spectrum for Pd/Au on 
p-GaN, 900°C, 15 seconds in N 2 , showing Au, Pd, Ga and N 
present in this area. 







4e-4 - 

2e-4 - 

Oe+0 - 

-2e-4 - 

-4e-4 - 


-2 2 
Voltage (V) 

Figure 5.11. I-V curves for sputter deposited Au on 
elecron beam evaporated Ni (2400/500A) contacts to p-GaN, 
heat treated in forming gas. 






3e-4 - 

2e-4 - 

le-4 - 

0e+0 - 

-le-4 - 

-2e-4 - 



Figure 5.12. I-V curves for Ni/Au contacts to p-GaN with 
varying layer thicknesses. 



8e+4 - 

2e+4 - 


•H ^ 

co 6e+4 - 



m 4e+4 - 

Oe+o ktJbamaMt**k 



10 15 20 

Sputter Time (Min) 


Figure 5.13. AES depth profile for Ni/Au on p-GaN, as- 



4-j 2e+4 





le+4 - 


2 3 4 5 6 
Sputter Time (Min) 

Figure 5.14. AES depth profile for Ni/Au (500/1000A) on 
p-GaN, heat treated at 400°C, 5 minutes in N 2 . 



2.0e+4 -j ■ 

1.5e+4 - 






H 1.0e+4 - 




Sputter Time (Min) 

Figure 5.15. AES depth profile for Ni/Au on p-GaN, heat 
treated at 600°C, 30 minutes in forming gas. 







2e-4 - 

le-4 - 

5e-5 - 

0e+0 - 

-5e-5 - 

-le-4 - 

-2e-4 - 




Ni/C/Au as-depbsited 

Ni/C/Au 200°C,i 5 min 

■ Ni/C/Au 400°C,;5 min 
O Ni/Aii as-deposited 

^....lli/Aii..2.0.0. o C,,..5:.min 

D Ni/Avj 400 C C, 5; min 


-2 2 

Voltage (V) 

Figure 5.16. I-V curves for Ni/Au and Ni/C/Au on p-GaN, 



2.0e+4 - 


"S 1.5e+4 - 


H 1.0e+4 - 

5.0e+3 - 



M/ 3t rJT.^=' "katV 

10 15 20 

Sputter Time (Min) 


Figure 5.17. AES depth profile for Ni/C/Au on p-GaN, 






m 1.5e+4- 


m 1.0e+4 - 

5.0e+3 - 


„..,.. Au 

k'ir i M.yww 

Sputter Time (Min) 

Figure 5.18. AES depth profile for Ni/C/Au on p-GaN, heat 
treated at 400°C for 30 minutes in forming gas. 



le+6 - 

4 6 8 10 

Sputter Time (Min) 


Figure 5.19. SIMS depth profile for Ni/C/Au on p-GaN, 


le+3 - 
le+2 - 
le+1 - 

k i-^V/^ Ga 






Sputter Time (Min) 



Figure 5.20. SIMS depth profile for Ni/C/Au on p-GaN, heat 
treated at 400°C, 30 minutes in forming gas. 




le-2 - 

5e-3 - 

Oe+0 - 

-5e-3 - 

-le-2 - 




• As-Deposited 

° 200°C, 5 min 

■ 400°C, 5 min 

D 900°C, 15 sec 

■2 2 

Voltage (V) 



Figure 5.21. I-V data for Cr/Au on p-GaN, 




•h l.5e+4 - 



G 1.0e+4 - 

5.0e+3 - 


0.0e+0 Uw» r»<-«- « 

10 15 20 

Sputter Time (Min) 


Figure 5.22. AES depth profile for Cr/Au on p-GaN, 



; > .Au 

2.0e+4 - 



1.5e+4 -i 


l.Oe+4 - 
5.0e+3 - 


fcni« < x'*-fc«an 

10 15 20 25 

Sputter Time (Min) 

Figure 5.23. AES depth profile for Cr/Au on p-GaN, heat 
treated at 200°C for 5 minutes in N 2 . 


2.0e+4 - 






Sputter Time (Min) 


Figure 5.24. AES depth profile for Cr/Au on p-GaN, heat 
treated at 400°C for 5 minutes in N 2 . 


1.2e+4 - 


■H 8.0e+3 - 




4.0e+3 - 


^'^^ W ^ A 

,.0e+0 ^Uifcirii^gg!^g^ 

••.•:•■.!.•• —ii'st »■•• ,. ,,f. 



10 20 

Sputter Time (Min) 


Figure 5.25. AES depth profile for Cr/Au on p-GaN, heat 
treated at 900°C for 15 seconds in N . 


♦ A 

l i *> 




o ®^sa o 

Figure 5.26. Low magnification XTEM (x73K) of Cr/Au on 
p-GaN, as-deposited, showing planar metal layers. 


Figure 5.27. SADP from an as-deposited Cr/Au contact 







5.0 k#v 10.0 










Figure 5.28. EDS data for the regions indicated from Figure 





Hole where GaN 
milled away 

Figure 5.29. Low magnification (x20K) XTEM for Cr/Au on 
p-GaN heat treated at 900°C for 15 seconds in N 2 . 







Figure 5.30. EDS data for Cr/Au on p-GaN, 900°C, 15 seconds 
progressing across the metal layers from the bulk to the 


Figure 5.31. High resolution XTEM of Cr/Au on p-GaN heat 
treated at 900 C C for 15 seconds showing Ga surface 






5 . O 


io . o 

Au:Cr region 

Ga-rich region 

Figure 5.32. EDS data for Cr/Au on p-GaN, 900°C, 15 seconds 
for the micrograph in Figuure 5.31. 


Figure 5.33. SADP from Cr/Au on p-GaN heat treated at 
900°C for 15 seconds showing Au and Cr. 


Figure 5.34 . SADP from Cr/Au on p-GaN heat treated at 
900°C for 15 seconds showing Ga-rich phase. 


4.E-04 T 


o°x x 

o x 
o v x 



v x 

•* n D «»* X X U 

° » X*0 

° B * xV 
♦ v x o 

X ° 

* V 

X O 



2 4 

-4.E-04 1 
Voltage (V) 


80 K 

140 K 


200 K 


260 K 


320 K 

380 K 

Figure 5.35 . Temperature dependent I-V for Ni/Au on p-GaN, 
heat treated at 600°C for 5 minutes. 



Figure 5.36 . Temperature dependent I-V for Pd/Au on p-GaN, 
heat treated at 600°C for 5 minutes. 





3.E-04 T 






\bltage (V) 









♦ xo 




140 K 




260 K 





Figure 5.37. Temperature dependent I-V for Cr/Au on p-GaN, 
heat treated at 600°C for 5 minutes. 







423 nm 
2.93 eV 

900°C, 15 seconds N 2 
As -grown 

I 1 1 1 — 

350 400 450 500 550 

Wavelength (run) 



Figure 5.38. PL data for p-GaN sample as-grown and 
following a 900°C, 15 second heat treatment showing levels 
from Mg complexes. 






900°C, 15 seconds N 2 
As -grown 

372 nm 
3.33 eV 

360 nm 
3.44 eV 





365 370 375 

Wavelength (nm) 


Figure 5.39. PL data for p-GaN sample as-grown and 
following a 900°C, 15 second heat treatment showing energy 
region near the bandgap. 


— ♦— lVNi/Au 



— X— IV W/Au 

-G— 3V W/Au 

-0-5V W/Au 

— t— IV Cr/Au 

-H-3V Cr/Au 

—•—5V Cr/Au 

28 : xxx-x-x-x-x-x— x— I 





Figure 5.40. Plot of ln(Io/T 2 ) vs 1/T from 100-400 K for 
contacts to p-GaN. 

-19 1 


-20 j 
-211 = 








-28 1- 

-29 J 




0.0025 0.0026 0.0027 0.0028 



0.003 0.0031 0.0032 0.0033 

Figure 5.41. Plot of ln(Io/T A 2) vs 1/T from 300-400 K for 
contacts to p-GaN. 



-7 ■ 



-«— 200KNi/fou 

-§— 30CK Ni/Au 

-*-40CK Ni/Au 

-X-20CK W/Au 


-e-30CK RJ/Au 

-B— 40CK W/Au 

— f-20CKCr^u 


-»— 40CKCr/fcu 

2 3 

Figure 5.42. Plot of In (I) vs V from 0-5 V for contacts to 



-4-200K Ni/fu 
-A-400K NiyJu 
-X-20OC Vd/fu 
— f— 200KCc/fti 
-X-30CK Cc/Ai 




Figure 5.43. Plot of In (I) vs V from 3-5 V for contacts to 


p-ZnTe Contacts 

The formation and characterization of sputter deposited 
Au contacts to nitrogen doped p-ZnTe was demonstrated. 
Based on current-voltage data, the electrical properties of 
the contacts were rectifying as-deposited with a breakdown 
voltage of 0.5 V. These values are close to the value of 
0.66 eV predicted using the Schottky model which strongly 
suggests the Fermi level was not pinned in ZnTe. 

Upon heating at T*200°C, ohmic behavior was observed 
from the I-V data. Through the use of AES and SIMS depth 
profiles, this ohmic behavior was credited to the creation 
of near-surface acceptors due to doping of Au (Ev=+272 meV) 
without inter facial phase formation. This led to increased 
conductivity through the near surface layer, presumably by 
carrier hopping between deep levels in the lattice. A 
maximum current density of 2.3 A/cm 2 at 5 V was obtained 
following a 250°C, 15 minute heat treatment in forming gas. 
At 350°C diffusion of Au was extensive and resulted in 
compound formation, proposed to be AuTe 2 , which led to 
severe degradation of the contact properties. 


Thus for Au/p-ZnTe contacts, diffusion of Au at T*250°C 
formed ohmic contacts, while at 350°C, high resistivity- 
contacts were obtained due to widespread Au diffusion and 
c ompound f o rma t i on . 

p-GaN Contacts 

For contacts to p-GaN, the large band gap (E g =3.4 eV) 
and electron affinity (x=3. 3-4.1 eV) required interfacial 
reactions to take place for an ohmic contact to be formed. 
The best contact schemes to p-GaN took advantage of 
dissociation of the GaN to allow for increased doping of the 
semiconductor near-surface region. Au contacts were sputter 
deposited and Pd/Au, Ni/Au, Ni/C/Au and Cr/Au were deposited 
by electron beam evaporation onto Mg doped p-GaN. 

I-V data showed rectifying contacts for all metal 
schemes as-deposited. Pure Au contacts showed no reaction 
with GaN up to 400°C and resulted in rectifying behavior. 
Pd/Au contacts were rectifying under all process conditions 
with increased conduction attributed to breakup of the 
interfacial oxide below 400°C and dissociation of the GaN 
following a 900°C, 15 second RTA heat treatment, with Pd or 
C potentially acting as an acceptor. 

Insertion of a Ni layer to form a Ni/Au contact led to 
nearly linear I-V curves and dissociation of the GaN as 
evidenced by AES depth profiles at 600°C. The increased 

conductivity was attributed to increased doping of the GaN 
near-surface region by interfacial carbon or Ni acting as an 
acceptor in the GaN lattice. Introduction of an 
intermediate C layer to form a Ni/C/Au contact did not 
decrease the contact resistance. 

Films of Cr/Au provided the best ohmic contacts with 
P c i4.3xl0" 1 Qcm 2 following a 900°C, 15 second RTA. At this 
temperature Cr caused the dissociation of GaN and was 
postulated to act as a substitutional acceptor sitting on 
the Ga sublattice. GaN doping was postulated to occur by an 
interstitialcy diffusion mechanism based on similar data for 
Cr in GaAs. This led to significant reduction of the 
depletion layer thickness and better carrier transport. 

Temperature dependent I-V measurements showed Ni/Au, 
Pd/Au and Cr/Au to all have components of both thermionic 
and field emission, with Cr being dominated at RT by field 
emission. This is consistent with the dissociation of GaN 
by Cr and enhanced diffusion to create increased doping in 
the near-surface region. 


D-ZnTe Contacts 

For contacts to p-ZnTe, work must be continued to find 
a contact scheme that provides low resistance thermally 
stable contacts. Au has been shown to form ohmic contacts 
to p-ZnTe, but the contact resistance is too high for 
practical devices. Unfortunately, there will probably not 
be a large interest in continued work with this 
semiconductor due to the move within the industry away from 
II-VI materials and towards III-V materials. 

p-GaN Contacts 

The advent of blue laser diodes made from GaN requires 
that the study of ohmic contacts to the p-type material 
continue until a suitable low resistance contact is found. 
One possible contact scheme for these low resistance 
contacts will include a doping layer, such as Mg, in the 
contact scheme in an effort to increase the near-surface 
carrier concentration. If this method does not prove 
feasible, p-GaN contacts will most likely need to be formed 


using an MBE process, similar to that used in ZnSe. It is 
believed that a simple metal contact scheme will not be able 
to provide the low resistance contacts needed for practical 
devices, and thus the contact scheme will have to consist of 
a graded region between the GaN and a semiconductor that can 
be very highly doped p-type, for example GaAs. 


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Jeffrey Todd Trexler was born on October 15, 1970, in 
Allentown, PA, and was raised in Bethlehem, PA. He 
graduated from Freedom High School in 1988. Upon completing 
high school he enrolled at Clemson University in Clemson, 
SC, and majored in ceramic engineering. In May of 1992, he 
graduated cum laude from Clemson University with a 
bachelor's degree. In August of 1992, Jeffrey began 
graduate school at the University of Florida in Gainesville, 
FL, majoring in materials science and engineering. Under 
the supervision of Dr. Paul Holloway, Jeffrey received his 
master's degree in materials science and engineering in May 
1996 and his PhD in materials science and engineering in 
December 1997. 


I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, as 
a dissertation for the degree of Doctor of Philosophy. 

Paul H. Holloway, Chaii 
Professor of Materials 
Science and Engineering 

I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, as 
a dissertation for the degree of Doctor of Philosophy. 


Cammy R. Abernathi 
Professor of Materials 
Science and Engineering 

I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, as 
a dissertation for the degree of Doctor of Philosophy. 

Kevin S. Jones- 
Associate professor of 
Materials Science and 

I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, as 
a dissertation for the degree of Doctor of Philosophy. 

Professor of Materials 
Science and Engineering 

I certify that I have read this study and that in my 
opinion it conforms to acceptable standards of scholarly 
presentation and is fully adequate, in scope and quality, as 
a dissertation for the degree of Doctor of Philosophy. 

Timothy' J.I Anderson 
Professor of Chemical 

This dissertation was submitted to the Graduate Faculty 
of the College of Engineering and to the Graduate School and 
was accepted as partial fulfillment of the requirements for 
the degree of Doctor of Philosophy. 

December, 1997 


Winfred M. Phillips 
Dean, College of 

Karen A. Holbrook 
Dean, Graduate School 





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